Advanced Structured Materials
For further volumes: http://www.springer.com/series/8611
Vikas Mittal Jin Kuk Kim Kaushik Pal Editors
Recent Advances in Elastomeric Nanocomposites
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Editors Dr. Vikas Mittal Berner Weg 26 67069 Ludwigshafen Germany e-mail:
[email protected] Present Address Dr. Vikas Mittal Chemical Engineering Program The Petroleum Institute Abu Dhabi 2533 UAE e-mail:
[email protected] Prof. Dr. Jin Kuk Kim Department of Polymer Science and Engineering School of Nano and Advanced Materials Engineering Gyeongsang National University Gazwa-Dong 900, Jinju, Gyeongnam 600-701 Korea, Republic of (South Korea) e-mail:
[email protected] Dr. Kaushik Pal Department of Polymer Science and Engineering School of Nano and Advanced Materials Engineering Gyeongsang National University Gazwa-Dong 900, Jinju, Gyeongnam 600-701 Korea, Republic of (South Korea) e-mail:
[email protected] ISSN 1869-8433
e-ISSN 1869-8441
ISBN 978-3-642-15786-8
e-ISBN 978-3-642-15787-5
DOI 10.1007/978-3-642-15787-5 Springer Heidelberg Dordrecht London New York Ó Springer-Verlag Berlin Heidelberg 2011 This work is subject to copyright. All rights are reserved, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilm or in any other way, and storage in data banks. Duplication of this publication or parts thereof is permitted only under the provisions of the German Copyright Law of September 9, 1965, in its current version, and permission for use must always be obtained from Springer. Violations are liable to prosecution under the German Copyright Law. The use of general descriptive names, registered names, trademarks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. Cover design: WMXDesign GmbH, Heidelberg Printed on acid-free paper Springer is part of Springer Science+Business Media (www.springer.com)
Preface
Elastomers are a very important class of polymer materials and the generation of their nanocomposites by the incorporation of nano-fillers has led to the enhancement of their properties significantly and hence expansion of their application potential. The book aims to specifically review the recent progresses in the synthesis, processing as well as applications of the elastomeric nanocomposites. The contents of the book are classified into three broad categories: first one dealing with introduction and preparation of the elastomeric nanocomposites, the second one focusing on the characterization and properties of the formed composites, whereas the third one describing the application potential of these materials. Chapter 1 describes the role of different nanoparticles in reinforcing elastomers. Homogenous dispersion of the filler and subsequent microstructure development in the composites have been focused upon. Chapter 2 explains the important synthesis methodology of in situ preparation of elastomeric nanocomposites. Other synthesis methodologies have also been described in brief. Chapter 3 focuses on the relaxation phenomena in elastomeric nanocomposites. The results are presented for both non-polar and polar polymer matrices. Modeling and simulation of nanocomposite processing have been described in Chap. 4 using molecular dynamics and Monte Carlo methodologies. Chapter 5 shows the deformation induced structural changes in elastomeric nanocomposites. Polymer matrices reinforced with various fillers like clay, nanofibers, nanotubes, carbon black, etc. have been considered. Thermally stable and flame retardant elastomeric nanocomposites are the focus of Chap. 6 whereas recycling of the elastomeric nanocomposites has been demonstrated in Chap. 7. In the applications section, Chap. 8 describes the considerations for the use of elastomeric nanocomposites in tyre applications. Subsequently, synthesis of nanocomposites suitable for use in tyre applications has been reported. Chapter 9 shows the use of elastomeric nanocomposites made from polyurethane and epoxy matrices for potential use in packaging applications. Elastomeric nanocomposites suitable for biomedical applications have been described in Chap. 11. Considerations of elastomeric nanocomposite systems for use in aerospace applications have been focused in Chap. 12. Chapter 13 describes the friction and wear of polymer nanocomposites v
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containing clay and nanotubes as fillers thus providing the potential of their use in such applications. The editors would also like to express heartfelt thanks to Springer, Germany for their acceptance to publish the book. Dr. V. Mittal Prof. Dr. J.K. Kim Dr. K. Pal
Contents
Part I
Introduction & Preparation
Role of Different Nanoparticles in Elastomeric Nanocomposites . . . . . Jin Kuk Kim, Kaushik Pal and V. Sridhar
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In Situ Synthesis of Rubber Nanocomposites . . . . . . . . . . . . . . . . . . . Massimo Messori
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Part II
Characterization & Properties
Relaxation Phenomena in Elastomeric Nanocomposites. . . . . . . . . . . . G. C. Psarras and K. G. Gatos
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Modeling and Simulation of Polymeric Nanocomposite Processing . . . Teik-Cheng Lim
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Deformation-Induced Structure Changes in Elastomeric Nanocomposites. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Shigeyuki Toki and Benjamin S. Hsiao
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Thermally Stable and Flame Retardant Elastomeric Nanocomposites. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . O. Cerin, G. Fontaine, S. Duquesne and S. Bourbigot
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Recycling of Elastomeric Nanocomposites . . . . . . . . . . . . . . . . . . . . . L. Reijnders
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Part III
Contents
Application
Elastomeric Nanocomposites for Tyre Applications. . . . . . . . . . . . . . . Kaushik Pal, Samir K. Pal, Chapal K. Das and Jin Kuk Kim
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Elastomer Clay Nanocomposites for Packaging. . . . . . . . . . . . . . . . . . V. Mittal
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Elastomeric Nanocomposites for Biomedical Applications . . . . . . . . . . Nicole Fong, Anne Simmons and Laura Poole-Warren
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Actuators and Energy Harvesters based on Electrostrictive Elastomeric Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Kaori Yuse, Pierre-Jean Cottinet and Daniel Guyomar Elastomeric Nanocomposites for Aerospace Applications . . . . . . . . . . James Njuguna, Krzysztof Pielichowski and Agnieszka Leszczyn´ska
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Friction and Wear of Rubber Nanocomposites Containing Layered Silicates and Carbon Nanotubes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . D. Felhös and J. Karger-Kocsis
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Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
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Part I
Introduction & Preparation
Role of Different Nanoparticles in Elastomeric Nanocomposites Jin Kuk Kim, Kaushik Pal and V. Sridhar
Abstract The use of different types of fillers in preparation and characterization of rubber based nanocomposites is depicted in this chapter followed by different types of techniques used for dispersion of nanofillers in the rubber matrix. The effects of the different fillers (CNT, CNF, nanographite) are generally discussed in terms of morphology (investigated by electron microscopy, small angle X-ray scattering, AFM), mechanical properties (modulus, stress and strain) and electrical properties (impedance, percolation threshold etc.). The main aim of this research is the dispersion of the carbonaceous materials in the elastomers and TPE gels which is a significant obstacle to the scientists. In this chapter we have tried to overcome this problem by using different types of fillers and it has been proved by the results obtained from different types of characterization methods.
1 Introduction Elastomeric composites are widely used because of their light weight, design flexibility, and processability. However, these composites exhibit less attractive mechanical properties such as low strength and low elastic modulus as compared to metals and ceramics. Adding micron- or nano-sized inorganic filler particles to
J. K. Kim (&), K. Pal and V. Sridhar Elastomer Lab, Polymer Science and Engineering, School of Nano and Advanced Materials, Gyeongsang National University, Jinju 660-701, Gyeongnam, Korea e-mail:
[email protected] K. Pal e-mail:
[email protected] V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_1, Springer-Verlag Berlin Heidelberg 2011
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reinforce the polymeric materials has been standard practice in the composite industry for decades. The term ‘‘filler’’ in rubber technology is often misleading, implying a material that is primarily intended to reduce the cost of the more costly rubber. But the modern-day fillers change one or more of these properties: optical properties and color; improve surface characteristics and dimensional stability; change thermal, magnetic, and electrical properties; improve mechanical properties, durability, and rheology; affect chemical reactivity, biodegradability, etc. The mechanical and physical properties of the composite are mostly dominated by the nature of the filler, whereas the polymer matrix determines the environmental characteristics of the composite. Therefore, the overall composite properties can be tailored to fit the desired application through the proper choice of filler and matrix resin. Therefore, a judicious choice of the type of filler and its concentration in the composite will augment the overall performance of the composite [1].
1.1 Classification of Fillers in Elastomers Although many types of fillers (e.g., fibers, whiskers, particulates, etc.) are widely used in the rubber industry, particulate fillers form a major share. Particulate fillers are broadly classified as reinforcing and non-reinforcing, depending on whether or not they enhance the performance characteristics of the final product.
1.1.1 Non-Reinforcing Fillers Fillers that only lead to small increases in viscosity of the compound, cause deterioration of the mechanical properties of the vulcanizate, and do not exhibit any reinforcing action are called non-reinforcing or inactive fillers (e.g., calcium silicate, chalk powder, etc.). These are often called extenders, and are used to reduce the production cost of rubber goods [2]. Generally, powder minerals are used as fillers and a few are listed below. The most widely used non-reinforcing fillers are china clay and calcium carbonate. Clay is probably the most commonly used non-reinforcing fillers and is classified as a hard clay or a soft clay, depending on their particle size (soft clay, [2 lm; and hard clay, \2 lm). These are low-cost fillers that can be used at high volumes to provide cheap compounds. However, clays impart improvement in properties such as hardness and abrasion resistance, and also reduce mold shrinkage. Calcium carbonate, also known as whiting, is extensively used in rubber formulations, primarily to impart color and reduce the cost the product [3, 4]. The two main types are ground limestone and precipitated calcium carbonate. The former is made by grinding mineral limestone and the latter is obtained by chemical precipitation from salt solution. Whiting provides smooth compounds that are easily processed and often used in microcellular foams. Ground whiting gives low tear
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resistance but the finer particle size precipitated materials exhibit better characteristics in this regard and have proved to have good hot tear strength. Other non-reinforcing fillers are barytes (barium sulfate), mica, titanium dioxide (TiO2), and silicates of calcium and zinc [5–9]. Barytes is generally used in medical applications because of their unique chemical resistance and inertness. Mica is used in composites where high thermal expansion and high resistivity are needed. Titanium dioxide is incorporated in rubber to mask the inherent color of the matrix.
1.1.2 Reinforcing Fillers The basic feature that distinguishes between non-reinforcing and reinforcing fillers is enhancement in performance characteristics such as the tensile strength, modulus, etc. of the composite. Reinforcement in vulcanized elastomers can be defined as the simultaneous increase in stiffness and resistance to fracture by the addition of filler [10–12]. A broader definition of reinforcement vis-à-vis the rubber industry is ‘‘the improvement in abrasion, tear, cutting and rupture resistance, stiffness and hardness of vulcanized compounds’’ [13]. Particulates such as carbon blacks and silica are the most widely used reinforcing fillers in the rubber industry. Carbon black, whose reinforcing character has been extensively exploited in rubber engineering, was discovered in the early twentieth century (at Silverton, United Kingdom, in 1910) and is the most widely used particulate filler in the rubber industry. Carbon black reinforcement became a subject of scientific interest only in the 1940s due to the development of suitable investigating tools and the growing use of synthetic rubbers in demanding applications, namely automotive and truck tires [14]. Carbon black is manufactured by a variety of processes, namely furnace, thermal, acetylene, and channel. The characteristics of the carbon blacks and their reinforcing nature in a polymer matrix are widely dependent on the type of manufacturing process.
1.1.3 Nano-Sized Fillers New-generation nano scaled fillers are challenging the domination of traditional fillers such as carbon blacks and silica in the rubber industry. Nanoscaled fillers such as layered silicates, carbon nanotubes, carbon nanofibers (CNFs), exfoliated graphite, etc. dispersed as a reinforcing phase in an elastomer matrix are emerging as a relatively new form of useful materials. These composites exhibit a change in composition and structure over a nanometer length scale and possess remarkable property enhancements relative to the pure polymer. Owing to the nanometer-size particles obtained by dispersion, these nanocomposites exhibit superior mechanical, thermal, optical, and dynamic mechanical properties at lower concentrations compared with either the pure polymer or conventional micron-sized composites. Their unique properties stem from a combination of factors such as their high
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aspect ratio (length to diameter), surface area, and the molecular bonds formed between the nano-sized filler and the polymer during compounding. The utility of nano-scaled fillers as possible reinforcing agents in elastomeric composites can be traced to the earliest work of Schmidt, who reported that colloidal stannic oxide, silica, Prussian Blue, polystyrene, and casein have very strong reinforcing effects in styrene butadiene rubber (SBR) and concluded that the ‘‘small particle size of the pigment (filler) is of prime importance in elastomer reinforcement, whereas the chemical nature of the filler appears to be of secondary importance’’ [15]. Although this hypothesis cannot be conclusively proven experimentally, there is, nevertheless, good evidence to show that small particle size is a necessary requirement, and very likely the predominant requirement for reinforcement effect in rubber is the small size of the fillers. Consequently, one can observe an increase in reinforcement with decreasing primary particle size in carbon blacks as reported by Studebaker [16], who showed a gradation of properties from the thermal blacks, with particle diameters above 300 nm, that show little reinforcement, through the ‘‘semi-reinforcing’’ furnace grades, with particle diameters in the region of 100–200 nm, thence to the ‘‘high abrasion’’ furnace grades (about 40 nm) and finally to the ‘‘intermediate super-abrasion’’ or ‘‘super-abrasion’’ grades with primary particle diameters below 35 nm. Wagner [17] reported that nano-sized fillers such as soft clays having a particle in the range of 1,000–8,000 nm as ‘‘diluent’’ fillers and hard clays, zinc and titanium oxides precipitated calcium carbonates in the range of 100–1,000 nm as ‘‘semi-reinforcing’’ and precipitated calcium carbonates, silicas, calcium silicates or silicoaluminates, or anhydrous silicas with particle diameters in the range of 10–100 nm as ‘‘reinforcing’’ fillers in natural rubber. Theoretical investigations on the utility of nano-scaled fillers in elastomers have been carried out by Fowkes and Gent [18–20]. Fowkes was the first researcher to report that when functional filler particles such as carbon blacks are dispersed in a rubber matrix, the polymer wets and adheres to the surface and is held by moderate intermolecular attractive forces and by surface tension. Huber and Heinrich [21, 22] presented detailed theoretical investigations concerning the hydrodynamic reinforcement contribution in elastomeric composites with rigid filler particles of fractal nature (carbon black and silica aggregates), spherical core–shell particles with a soft core and hard shell, spherical core–shell particles with a hard core and soft shell (carbon black particles with bound rubber). In the context of carbonblack-filled elastomers, the contribution to reinforcement on small scales can be attributed to the complex structure of the branched filler aggregates, as well as to a strong surface polymer interaction, leading to the so-called bound rubber. Thus, the filler particles are coated with polymer chains, and the binding (physically or chemically) of elastomer chains to the surface of the filler particles significantly changes the elastic properties of the macroscopic material. On larger scales, the hydrodynamic aspect of the reinforcement dominates the physical picture. Hydrodynamic reinforcement of elastic systems plays a major role not only in carbon-black-filled elastomers, but also in composite systems with hard and soft
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inclusions. Finally, at macroscopic length scales, the existence of filler networking at medium and high filler volume fractions plays the dominate role. 1.1.4 Carbonaceous Nano-Fillers Carbon black is the most commonly used filler in elastomers. The past few years have seen the extensive use of nanoparticles because of the small size of the filler and the corresponding increase in the surface area, allowing achieving tremendous increase in mechanical properties even at very low filler loadings. Recently, considerable research efforts have focused on nano-scale variants of carbon black (viz. carbon nanotubes, carbon nanofibers, and exfoliated nanographite) as possible reinforcing fillers in elastomers. Among these, nanotubes are attracting the most attention. Carbon Nanotubes The discovery of carbon nanotubes can be traced back to the origin of fullerene chemistry (Buckyball, C60) in 1985 [23]. In 1991, Ijima [24] discovered carbon nanotubes (CNTs) that are elongated fullerenes, where the walls of the tubes are hexagonal carbon and often capped at each end. Carbon nanotubes are needleshaped, single crystals whose properties depend on the atomic arrangement, chirality, diameter, and length of the tube and the overall morphology. Carbon nanotubes can be synthesized by different techniques, including arc-discharge laser ablation and various catalytic chemical vapor depositions (CCVDs) [25–29]. There are two types of CNTs: (1) multiwalled carbon nanotubes (MWCNTs) and singlewalled carbon nanotubes (SWCNTs). An SWCNT (or SWNT) is best described as a twodimensional graphene sheet (a hexagonal array of carbon atoms) rolled into a tube with pentagonal rings as end caps. SWNTs have aspect ratios of 1,000 or more, and an approximate diameter of 1 nm, or 10 Å. Similarly, MWCNTs (or MWNTs) can be described as multiple layers of concentric graphene cylinders, also with pentagonal ring end caps. Both SWCNTs and MWCNTs have physical characteristics of solids and are microcrystals with high aspect ratios of 1,000 or more, although their diameter is close to molecular dimensions. Table 1 shows the theoretical and experimental properties of carbon nanotubes. From Table 1 it is clear that CNTs have a unique combination of mechanical, electrical, thermal, and electrical properties that make them excellent candidates to substitute or complement the conventional fillers in the fabrication of elastomeric composites. Carbon Nanofibers The history of carbon nanofibers goes back more than a century. In a patent published by Hughes and Chambers [30] in 1889, it was reported that carbon
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Table 1 Theoretical and experimental properties of carbon nanotubes Property CNT Graphite Specific gravity Modulus Strength Resistivity Thermal conductivity Specific surface area
3
0.8 g/cm for SWNT 1.8 g/cm3 for MWNT *1 TPa for SWNT *0.3–1 TPa for MWNT 50–500 GPa for SWNT 10–60 GPa for MWNT 5–50 lX-cm 3,000 W/m-K 100–250 m2/g
2.26 g/cm
3
Carbon black 1.2–2.4 g/cm3
1 TPa (in plane)
–
–
–
50 lX-cm (in plane) 3,000 W/m-K (in plane) 10–20 m2/g
25–40 lX-cm – 8–130 m2/g
filaments were grown from carbon-containing gases using a metallic crucible. However, 80 years later it was Robertson [31] who was the first to recognize that it was the interaction of methane and metal surfaces that led to the growth of graphitic carbon at relatively low temperatures. For the first 80 years of this century, the occurrence of carbon nanofibers—then often referred to as carbon filaments or filamentous carbon—was considered a nuisance. In the 1980s, several workers explored the use of carbon nanofibers for such applications as reinforcing fillers in polymers and as catalyst support material. Carbon nanofibers based on different precursors (PAN, rayon, and pitch) show various kinds of surface structures [32]. The basic structure of a typical carbon fiber (PAN) consists of long primary units (lateral aromatic molecules) lying parallel to the fiber axis and bonding together to form a stretched network of branched fibrils that apparently run the full length of the fiber. They form ribbon-shaped monatomic layers of sp2-type carbon, with an average width of 6 nm and a length of several hundred nanometers [33, 34]. The ribbons display an irregular contour, which appears ‘‘turbostratic’’ in the crystalline structure, and a certain number of these ribbons run parallel to form the microfibrils that have a preferred orientation parallel to the fiber axis.
Layered Nano Dimensional Fillers Layered fillers such as mica and talc have been used as fillers for elastomers due to their low cost, easy availability and outstanding electrical, heat, and chemical resistance. The extent of reinforcement by mica depends on the aspect ratio (i.e., the, ratio of length to thickness). The breakdown of mica particles during mixing and processing causes a reduction in the aspect ratio, which can be countered partially by using special equipment such as hot runners, longer sprues, and streamlining to reduce particle breakdown [35]. But mica is an inert filler devoid of any chemical groups, thereby having very low reinforcing capability when compared to traditional rubber fillers such as carbon black and silica. A variety of surface treatments, such as reacting mica platelets with titania, chlorinated
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paraffins, and silane coupling agents, has been studied in order to increase the mica-polymer interactions [36–38]. Other physical means, such as ultrasonic delamination, have been investigated by Tausz and Chaffey [39]. Mica-enhanced anisometric swelling of nitrile rubber vulcanizate and reduction of permeability of gases have been reported [40]. Mica, due to its layered structure, has good shock-absorbing and vibration-damping properties and has been investigated in polyurethane [41]. The anisotropic nature of mica will be useful in barrier applications [42]. Yoshikawa et al. [43] patented a rubber hose based on nitrile rubber–mica composites for fuel line systems in which swelling, permeability, and low-temperature properties were optimized by selection of the mica filler without any surface treatment. Debnath et al. [44] reported the effect of a silane coupling agent on vulcanization conditions, network structure, polymer-filler interaction, physical properties, and failure mode of mica-filled styrene butadiene rubber. The other type of layered filler is nanoclay. During the past 10 years, nanoclays have raised considerable interest in the scientific community due to their size and wide range of outstanding material properties. Of specific interest is the use of nanoclay-reinforced composites for structural applications, where experimental results demonstrate that substantial improvements in the mechanical behavior of the polymer can be attained through the addition of very small amounts of nanoclay. Understanding the structure–property relations in polymer–clay nanocomposites is of great importance in designing materials with the desired properties. In general, the improvement in properties of a polymer–clay nanocomposite originates from the nature of the layered inorganic fillers, their extent of dispersion in the polymer matrix, and by the distinctive interactions between a specific polymer and the nano-clay. Significant improvements in the physical and mechanical properties of layered alumina silicate reinforced polymeric composites have been reported. Some of the properties that undergo substantial improvement due to addition of nano-phases include mechanical properties such as strength, modulus, and dimensional stability; decreased permeability to gases, water, and hydrocarbon vapors; thermal stability and heat distortion temperature; flame retardance; surface appearance; and electrical conductivity. The most common types of nano-reinforcements are layered natural silicates such as montmorillonite, hectorite, and saponite [45–47]. Synthetic clays such as synthetic mica and synthetic hectorite have also been used [48, 49]. However, the most extensively used and studied is montmorillonite (MMT). Dispersion within the polymer matrix is required to form the inorganic–organic composite. To make MMT compatible with the organic polymer matrix, a cation exchange reaction is used to replace the sodium, potassium, and calcium ions with alkyl ammonium long-chain ions. This result in expansion between the clay galleries due to the larger molecules inserted between the layers. The reaction also changes the clay from hydrophilic to hydrophobic, making it more compatible with the organic matrix. Clay minerals are composed of nanoplatelets with a length of about 0.5–1 lm and a thickness of 1 nm, leading to large aspect ratios (500– 2,000).
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Processing of nanoclay composites is a real challenge because the elementary nanolayers or tactoids of a few platelets must be uniformly distributed in order to develop the highest amount of interfacial area, thus increasing the contact surface with the polymer chains. The main advantage of polymer-layered silicate nanocomposites is that great improvements in mechanical properties can be obtained even at very low clay content, which is economically interesting. However, to achieve this, there must be good dispersion of nanoclay platelets in the polymer matrix. But the intercalation or exfoliation of polymer chains in the nanoclay inter-gallery layers is a complex reaction that depends on many variables, such as polymer–clay interactions, chain conformation of the polymer, etc. However, the general consensus is that the major difficulty for the polymer chain to intercalate in the inter-gallery layers arises from the inherent structure of the clay particles, which have closely stacked silicate sheets. The diameter of these sheets typically lies between 20 and 200 nm, while the thickness is on the order of 1 nm. The spacing between the sheets is roughly on the order of 1 nm, which is less than the radius of gyration of typical polymers. Therefore, there is a large entropic barrier that inhibits the polymer from penetrating this gap and intermixing with the clay. Even when the sheets are successfully separated and interspersed into the polymer matrix, these high aspect-ratio platelets can form ordered or crystalline structures within the polymers or can phase-separate from the matrix material. To reduce the entropy of the nanoclay platelets, it is modified by treating with alkyl ammonium modifiers. These alkyl ammonium modifiers react electrokinetically with the cations inside the galleries of the clay by ion exchange, thereby causing the clay platelets to move apart, resulting in a reduction in the level of plateletplatelet attraction and in increased access for the polymer chains to intercalate [50]. Of the many techniques employed to achieve a high degree of dispersion of nanoclay platelets, in situ polymerization of monomers initially intercalated between silicate layers, and melt intercalation dispersing nanoclay platelets in a polymer solution are the most common [51–56]. There are few studies available on rubber-clay nanocomposites [57]. The important rubberclay nanocomposites thus far studied are elastomeric polyurethane, natural rubber, epoxidized natural rubber, styrene butadiene rubber, butadiene rubber, nitrile butadiene rubber, EPDM, silicone rubber, fluoroelastomers, butyl rubber, chlorobutyl rubber, etc. [58–68].
Nanographite Nanographite shows a similar layered structure and has a high aspect ratio in the range of 1,000–1,500 [69]. Graphite can provide additional advantages such as excellent electrical and thermal conductivity. Exfoliated graphite is prepared by rapidly heating a graphite intercalation compound (GIC). An exfoliated graphite nanoparticle is composed of stacks of nanosheets that can vary from 4 to 40 nm. They also show good affinity for both organic
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Table 2 Comparisons of properties of nanoclays and exfoliated nanographites Exfoliated clay Graphite nanoplatelets Physical structure Chemical structure Type of interactions between platelets Tensile modulus Tensile strength Electrical resistivity
Platelets, *1 nm 9 100–200 nm SiO2, Al2O3, MgO, K2O, Fe2O3 Hydrogen bonds, dipoles
Platelets, *1 nm 9 100–200 nm Graphene p–p
0.17 TPa *1 GPa 1010–1016 X-cm
Thermal conductivity Specific gravity
6.7 9 10-1 W/m-K
*1 TPa *10–20 GPa 50 9 10-6 X-cmk *1X-cm\ 3,000 W/m-Kk 6 W/m-K\ 2 g/cm3
2.8–3 g/cm3
compounds and polymers, therefore some monomers and polymers can be absorbed into the pores and galleries of exfoliated graphite. A comparison between layered silicates and exfoliated nanographite is shown in Table 2. From this table it can be observed that nanographite has additional advantages, including increased electrical and thermal conductivity. So, if one can effectively disperse nanographite in a polymer composite, many additional properties such as conductivity and electromagnetic shielding resistance can be obtained.
2 Problems Associated with Nanofillers in Elastomers 2.1 Dispersion It is recognized that there is still a long way to go to create CNT-reinforced composites fully realizing the potential of high stiffness and strength of CNTs. Many of the reported CNT polymer composites have not presented anticipated macroscopic properties such as high stiffness and strength, although the individual CNTs have excellent mechanical properties. There are many reasons that influence the macroscopic mechanical properties of nanocomposites, such as the dispersion, alignment, and waviness of CNTs. The dispersion of CNTs in polymers is relatively poor due to the nanotubes having a strong tendency to agglomerate as a result of their strong van der Waals’ forces. Their very stable chemical characteristics and lack of functional sites on the surface also complicate the dispersion issue. Moreover, the length of CNTs prepared from thermally activated CVD ranges from meters to several millimeters, which is undesirable for practical applications. Both physical and chemical approaches have been adopted to reduce the length of CNTs to a certain extent that is suitable for blending. The physical dispersion route generally includes ultrasonication, ball milling, grinding, and high-speed shearing. The as-prepared CNT product exists as loose
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multi-agglomerates, which can be separated by physical methods. Here we present a brief discussion regarding some of the problems associated with the fabrication of nano-scale filler-reinforced elastomers. There are four main system requirements for effective reinforcement [70]: (1) a large aspect ratio, (2) good dispersion, (3) alignment, and (4) interfacial stress transfer. The aspect ratio must be large to maximize the load transfer to the nanotubes. This is crucial in order to optimize composite strength and stiffness. Dispersion is probably a more fundamental issue. Dispersion of fillers in a polymer matrix results from the application of electrochemical and mechanical forces to the interface of the inorganic filler/polymer so as to cause complete de-agglomeration to the attired or original particle size in an organic phase, complete elimination of air voids and water, and the creation of a true continuous inorganic–organic composition [71]. The optimal performance of such nano filler reinforced polymer composites is achieved when the nano-sized fillers are uniformly dispersed in the polymer matrix. Fabricating such homogeneous mixtures poses considerable synthetic challenges. The major difficulty arises from the inherent anisotropic nature of the clay particles, which are composed of broad, closely stacked silicate sheets. The diameter of the sheets lies typically between 20 and 200 nm, depending on the specific type of clay, and each sheet is roughly 1 nm thick. The spacing (or ‘‘gallery’’) between the sheets is also on the order of 1 nm, which is smaller than the radius of gyration of typical polymers. Consequently, there is a large entropic barrier that inhibits the polymer from penetrating this gap and intermixing with the clay. Theoretical calculations made by Balazs’ group in a series of articles show that to understand the exfoliation process in polymer–clay nanocomposites, one must consider the kinetic and thermodynamic aspects of the penetration of polymer into the gap between the clay sheets [72–74]. So, based on this, the effective dispersion of nano-sized fillers in a polymer matrix remains more of an art than a precise science.
2.2 Low Mechanical Properties Nanotubes must be uniformly dispersed as isolated nanotubes individually coated with polymer. This is imperative in order to achieve efficient load transfer to the nanotube network. This also results in a more uniform stress distribution and minimizes the presence of stress-concentration centres. Alignment is, in some ways, a less crucial issue. From geometric considerations, the difference between random orientation and perfect alignment is a factor of five in composite modulus. While alignment is necessary to maximize strength and stiffness, it is not always beneficial. According to the Mori–Tanaka theory, aligned plate-like objects intrinsically reinforce less than aligned fibers [75–78]. Consequently, in the absence of disorder, modulus enhancement in clay nanocomposites would theoretically be less than that of carbon nanotubes. Angular averaging has less impact on platelets than on fibers,
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so randomly aligned platelets are ultimately more effective than randomly aligned fibers in terms of modulus enhancement [79]. The interface behavior can significantly affect the mechanical properties of nanocomposites. Nanotubes and nanofibers are normally produced at very high temperatures, thereby leading to the formation of highly graphitized carbon structures devoid of any surface groups. So, to increase the interactions, some sort of surface functionalization or activation of nanotubes and nanofibers is necessary. Rubber technologists were the first to recognize this phenomenon, and consequently a large amount of research and efforts has focused on this. They postulated that for a filler to be considered an effective reinforcing agent, primary valence bonding between polymer and filler is a prerequisite. The general conclusion of many studies is that higher the ‘‘surface activity’’ (due to presence of active surface groups such as carboxy, carbonyl, lactones, quinones, etc.) of the carbon black, the greater will be the reinforcement. While it is true, as indicated above, that reinforcement does occur in systems where strong interactions are highly unlikely, there is much evidence that a degree of reinforcement is less and is greater in fillers with strong interactions and is important in optimizing the properties of practical rubber compositions [80]. The most familiar example for understanding is the removal of reaction sites from carbon black (by graphitization) causes a drastic reduction in the modulus and abrasion resistance properties imparted by the parent (ungraphitized) carbon. Strong bonding may be of importance in at least two distinct ways. First, it may be used as a means of pulling apart filler agglomerates during the mixing process, thereby providing improved dispersion of the ultimate particles in the rubber, and opposing any tendency toward subsequent particle re-agglomeration [81]. Second, it may contribute to the adjustment of important physical characteristics, such as modulus, extensibility, and resilience in the vulcanizate. Carbon nanotubes are highly anisotropic in nature. Bradshaw et al. [82] proposed that curvature of embedded CNTs or ‘‘waviness’’ significantly reduced their reinforcement capabilities (by factors from 50 to 200) compared to straight CNTs. It was also noted that other indistinguishable factors contribute to the low values measured in experimental data, including weak interfacial bonding, insufficient dispersion, and degradation of the CNTs due to processing. At very low strains, it was suggested that the effect of poor interfacial shear strength should not affect the composite’s modulus, implying a measurable elastic response below the strain detrimental to the CNT interface [83]. Studies by Schadler et al. [84–86] have shown that carbon nanotubes in general do not bond well in polymers and their interactions result mainly from the weak Van der Waals forces. Consequently, CNTs may slide inside the matrix and may not provide much reinforcing effect. It is, however, important to assess whether the poor interface behavior is indeed responsible for the shortfall of CNT-reinforced composites in order to reach their expected properties. This is attributed to the absence of stress transfer to internal layers of MWNTs; only the outermost layer contributes to tensile reinforcement. It is noted that the additions of pristine and functionalized carbon nanotubes and carbon black particles lead to relatively large enhancements in Young’s modulus and fracture toughness [87, 88].
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3 Some Possible Solutions The high aspect ratio of the nanotubes coupled with a strong intrinsic van der Waals attraction between nanotubes combine to produce ropes and bundles of CNTs, particularly in SWNTs where the attractive force is approximately 0.5 eV per nanometer of nanotube-to-nanotube contact [89]. In the case of multi-walled carbon nanotubes (MWCNTs), the molecular forces between the individual MWCNTs might not be significant, so that sp2/sp3 orbit-hybridization in the walls of the MWCNTs rarely occurs. Therefore, MWCNTs are formed in the individual state with diameters at the high end of the nanometer scale. However, the dispersion of SWNTs is more difficult due to their highly crystalline nature and very strong sp2/sp3 hybridization.
3.1 Physical Techniques for Good Dispersion 3.1.1 Grinding and Rubbing There are few reports on the subject of rubbing or grinding carbon nanotubes to decrease the size. Rubbing is more destructive than any other method. The process introduces cuts and bends in the SWNT [90]. A less damaging method is chemically cutting SWNTs by grinding them in a fluid (alpha- or beta-cyclodextrin) using mortar and pestle. Both tube lengths and bundle diameters were noticeably reduced. Other grinding agents were used as well but did not give as good results; samples contained mostly long tubes [91].
3.1.2 High Shear Mixing-Shear Stress through a Nozzle Another mechanical method is to apply shear force to pull agglomerates apart— that is, high shear mixing. Usually narrow passages, and/or relatively high rates of flow, are required to generate high shear; in a lot of cases, rotor and stator construction is used [92]. Hilding et al. [93] applied a diesel fuel injector to create the high shear and correlated the reduction in viscosity with nanotube breakage.
3.1.3 High-Energy Ball Milling Pierard et al. [94] reported a comprehensive study on the effect of high-energy ball milling on the structure of SWNTs. Based on physical properties such as diameter, length, and surface area, they concluded that the optimum ball milling time is 2 h. Konya et al. [95] reported a simple ball milling technique in specific atmospheres to introduce functional groups such as thiol, amines, amides, carbonyl, etc.
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Lee et al. [96] reported a cryogenic crushing in the presence of liquid nitrogen to produce CNTs of 500-nm length. Other interesting technique to cut CNTs is the use of gamma-radiation in the presence of dilute sulfuric acid [97].
3.2 Chemical Techniques for Good Dispersion Carbon nanotubes and nanofibers are devoid of any chemical functional groups on their surfaces and hence have low interfacial interactions between them and the polymer matrix. Some researchers have made considerable efforts in the chemical modification of NTs, which might pave the way to many useful applications. The main approaches for the modification of these quasi-one-dimensional structures can be grouped into three categories [96]: (1) Covalent attachment of chemical groups through reactions with the ð-conjugated skeleton of CNT; (2) Non-covalent adsorption or wrapping of various functional molecules (3) Endohedral filling of their inner empty cavity A comprehensive review of the grafting and coating of CNTs by polymers is beyond the scope of this chapter. Some representative methods include attachment of long alkyl chains and polymers, fluorination, and radical reactions; these have provided access to tip and sidewall functionalization, eventually leading to a relative increase in solubility of CNTs in the polymer matrix [97–100]. The utility of silane coupling agents such as aminopropyl triethoxy silane and TESPD {poly(trimethyl dihydro quinoline)} to improve the rubber-filler interactions has been well investigated [101]. Shanmugharaj et al. [102] have functionalized CNTs with aminosilanes and explored the utility of the same as reinforcing filler in natural rubber. In addition to the above techniques, grafting or wrapping of conductive nano-sized metals on the surface of CNTs, especially for electromagnetic (EMI) shielding applications, have also been investigated. Techniques such as sol– gel, vapor deposition, and electrochemical methods have been used [103, 104]. The basic step for all chemical functionalizations of nanotubes or nanofibers is acid treatment. Although a concentrated H2SO4/HNO3 (vol:vol = 3:1) treatment is efficient in severing entangled nanotubes to enable their dispersion as individuals, damage to the tube-wall layers is serious and unavoidable; it has been reported that highly dispersed rod-like carbon can be achieved by sonication in negatively charged polyelectrolyte solutions [105, 106]. It has also been reported that acid treatment of CNT could improve the processability and performance of composites by introducing carboxylic acid groups on the surface of CNT, which leads to stabilization in polar solvents and helps to covalently link polymers. It must be mentioned that both acid treatment and ultrasonication have been reported to cause fragmentation (comminution) and damage of the nanotubes, which can be modeled using the moments of the length distribution. As with other solids, the breakage rate of carbon nanotubes depends on their lengths, with the
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longest particles experiencing the highest breakage rates. Perhaps the earliest reports of the effect of ultrasonic treatment on the properties of CNT was reported by Lu et al. [107], who reported extensive bending, buckling, damage, and formation of defects on the surface of nanotubes. The utility of surfactants to increase the dispersion of CNTs in the polymer matrix has been investigated. A surfactant’s property of accumulation at surfaces or interfaces has been widely utilized to promote stable dispersions of solids in polymeric media [108]. These amphiphilic molecules (i.e., compounds having both polar and apolar groups) adsorb at the interface between immiscible bulk phases, such as oil and water, air and water, or particles and solution, act to reduce the surface tension. The distinct structural feature of a surfactant originates from its ‘‘duality’’: the hydrophilic region of the molecule or the polar head group, and the hydrophobic region or the tail group that usually consists of one or few hydrocarbon chains. Although surfactants increase the dispersion of CNTs in the polymer matrix, it has many detrimental effects on the performance characteristics. Hernandez et al. have reported that by addition of the 1 wt% of CNT, which was initially dispersed by Triton X-100 surfactant (polyethylene oxide (9) nonyl phenyl ether), the improvement in storage modulus was only 50% when compared to over 230% in untreated CNT [109]. A similar tendency was found when the glass transition temperatures, Tg, of the neat polymer and surfactant-contained polymer were compared. This surfactant-induced plasticization effect was also reported for epoxy and polyethylene glycol matrices [110, 111].
3.3 Analysis of Dispersion In addition to the obvious difficulty in obtaining stable and homogeneous dispersions of nanotubes, another complication is finding a valid method to evaluate their state of dispersion. The most widely used visualization technique is microscopy: optical, scanning electron, and transmission electron microscopies. Visualization of CNT-based samples by optical microscopy enables to access mainly micrometer-sized agglomerates, while atomic force microscopy (AFM) is used to monitor suspended CNTs at the nano-scale level. But by using AFM, one can probe only a few nanotubes at a time, which might not be representative of the entire sample. Imaging of CNT-based polymeric composites by scanning or transmittance electron microscopy (SEM or TEM, respectively) often requires pretreatment by means of gold or carbon sputtering or microtome slicing of the sample, which might cause a defect in the original pattern of the composite. Solutions of carbon nanotubes are best viewed with cryo-TEM, which is ideally suited for imaging of wet samples. Scattering techniques such as x-ray diffraction, small angle neutron scattering, and dynamic light scattering have been used to study MWNT suspensions by particle size and agglomerate size, both in solution and polymer composites. The data analysis techniques have many complications, including assumptions of
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perfectly spherical particles, aligned rods, etc. As a result, the reported dimensions of nanotubes are in over-fitting range. The discovery of nanotube fluorescence has led to a precise method of detecting individual nanotube dispersion [112]. A nanotube in an aligned bundle does not emit because of energy transfer to neighboring tubes, particularly to the metallic ones. Thus, the dispersion process can be monitored by examining transient fluorescent emission as a function of various parameters, such as the type of surfactant used, sonication time, and surfactant concentration and functionalization [113]. In case of MWNTs, fluorescence spectra data can be misleading because MWNTs are a mixture of conducting (metallic) and semi-conducting nanotubes [114].
4 Utility of Carbonaceous Nanofillers in Elastomers and TPE Gels The next part of this chapter briefly discusses the preparation and properties of carbonaceous nano filler reinforced polymeric composites studied the following items by our group. • • • •
Exfoliated nano graphite in fluoroelastomers [115, 116] Multiwalled carbon nanotubes in fluoroelastomers [117] Vapor-grown carbon nanofibers in chlorobutyl elastomers [118] Multiwalled carbon nanotubes reinforced thermoplastic gels [119]
Elastomeric nanocomposites were prepared by traditional two-roll mill mixing technique, whereas swelling of SEBS thermoplastic in mineral oil was used to fabricate the TPE (thermoplastic elastomer) gels. Special attention focused on the dielectric relaxation of fluoroelastomer–nanographite composites, whereas straindependent dynamic mechanical properties were extensively studied for chlorobutyl-CNF composites.
4.1 Morphology The dispersion and morphology of the nanofillers in the composites have been studied by AFM, SEM, and TEM (Figs. 1, 2, 3).
4.1.1 Nanographite The dispersion of graphite platelets in fluoroelastomers has been studied using TEM andrepresentative microphotographs at 0.5 phr and 3.5 phr nanographite loadings are shown in Fig. 4a and b, respectively. From this figure it can be
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Fig. 1 General structure of sodium montmorillonite (NaMMT)
Fig. 2 General procedure used to make exfoliated nanographite
observed that the graphite nanosheets consist of thin graphite nano-lamellae with thicknesses of 1–5 nm or thinner, and with inter-gallery spacings of 3.37 Å. These inter-gallery distances reflect the XRD plots (Fig. 5), which show that both pure graphite platelets and graphite-FKM exhibit an intense peak at a diffraction angle of 26.4, corresponding to a basal spacing of 3.37 Å. A similar observation was reported by Yasmin and Daniel [120], who concluded that graphite platelets remain multilayered and maintain their original d-spacing. The dispersion of the nanographite platelets in the polymer matrix has been further studied by atomic force microscopy. Figures 6 and 7 show a representative topology of 0.5 and 3.5 phr nanographite reinforced. From these figures, multilayered nanographite platelets having a parallel orientation can be observed.
Role of Different Nanoparticles
Fig. 3 Dispersion and fabrication of nanoclay-reinforced polymer composites
Fig. 4 TEM of nanographite dispersed fluoroelastomer composites at a 0.5 phr and b 3.5 phr
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Fig. 5 WAXD diffractographs of the matrix and the nanographite/fluoroelastomer composites
Fig. 6 a Topology and 3-D AFM image of 0.5 phr nanographite-reinforced fluoroelastomer composites, respectively; b cropped topographical image of the nanographite platelets; and c line diagram
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Fig. 7 AFM topology images of 3.5 phr nanographite-reinforced fluoroelastomer composites
4.1.2 Carbon Nanofibers Sections of samples of chlorobutyl rubber compounds containing increasing amounts of VGCNFs (vapor-grown carbon nanofibers) were examined using SEM. Figure 8a–c show the SEM microphotographs at increasing VGCNF loadings. All the microphotographs show excellent distribution of the filler particles. From the figures it can be observed that at all filler loadings there is uniform distribution of the VGCNFs, the majority of which are individual nanofibers rather than aggregates, thus implying uniform distribution of nanofibers in the polymer matrix. SEM microphotographs (Fig. 8d, e) taken at lower length scales (2 lm) clearly show this phenomenon. The dispersion of VGCNFs in the polymer matrix was further studied by electron probe microanalysis (EPMA). The main advantage of EPMA is selective elemental scans. In the present study, carbon and oxygen scans have
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Fig. 8 SEM microphotographs of VGCNF-reinforced chlorobutyl elastomers at increasing filler loadings: a 5 phr, b 10 phr, and c 15 phr. SEM microphotographs of VGCNF-reinforced chlorobutyl elastomers at: d 5 phr, e 10 phr, and f 15 phr taken at smaller length scale (2 lm)
been carried out and a representative carbon and oxygen scan of a VGCNFloaded CIIR compound is shown in Fig. 9a and b, respectively. Figure 9a shows the distribution of nanofibers in the polymer matrix as the carbon scan, wherein it can be observed that VGCNFs are well dispersed with islands of carbon nanofibers, the majority of which are individual nanofibers rather than aggregates and implying uniform and excellent distribution of the nanofibers in the polymer matrix. Figure 9b shows the oxygen scan. Islands of dark color can be observed in the figures that indicate the presence of oxygen on the surface of the nanofibers.
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Fig. 9 Two-dimensional distribution of carbon nanofibers in chlorobutyl elastomer: a carbon scan and b oxygen scan
4.1.3 Carbon Nanotubes Figure 10a–c, and d show the SEM microphotographs of CNT-reinforced fluoroelastomers at increasing MWNT loadings. All the microphotographs show excellent distribution of the filler particles. From the figures it can be observed that at all filler loadings there is uniform distribution of the MWNTs, the majority of which are individual nanotubes rather than aggregates and implying complete exfoliation of the nanotubes in the polymer matrix. The MWNTs (bright dots) are found to disperse homogeneously in the rubbery matrix. The excellent distribution of nanotubes in the polymer matrix can be explained in terms of colloids, wherein the adsorbed polymer stabilizes the nanotube dispersion and protects it against bridging flocculation and depletion aggregation caused by the free polymer [121]. Representative TEMs of a 6 phr MWNT-reinforced fluoroelastomer composite are shown in Fig. 11a and b. Good dispersion of the nanotubes in the polymer matrix can be observed. From Fig. 11b, taken at a smaller length scale (20 nm), an absorbed polymer layer on the walls of nanotube can be observed, indicating wetting of the nanotubes in the polymer matrix. A novel technique, magnetic force microscopy (MFM), has been utilized to further study the dispersion of nanotubes in the polymer matrix. MFM, a variant of
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Fig. 10 Scanning electron microscopy of razor-cut section of MWNT-reinforced fluoroelastomer vulcanizates: a 1.5 phr, b 3 phr, c 4.5 phr, and d 6 phr (Reproduced from Ref. [117] with permisssion from Wiley Interscience)
Fig. 11 TEM microphotographs of MWNT-fluoroelastomer composites (Reproduced from Ref. [117] with permisssion from Wiley Interscience)
scanning force microscopy, is a tool capable of revealing a magnetic sample’s domain structure in real space. A representative topological picture of 6 phr nanotube reinforced fluoroelastomer is shown in Fig. 12. From the figure it is observed that the topographical image of MWNT/fluoropolymer composite is not as clear as has been observed by traditional techniques such as AFM and SPM; this can be attributed to the reduced ability of image contrast in MFM. It is generally known that it is not easy to separate the magnetic contrast from other background forces in MFM topography images. However, polymer/nanotube composites are two-phase materials with two distinct magnetic properties. The nanotubes are paramagnetic or diamagnetic, depending on their orientation, whereas the polymer
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Fig. 12 Topology of MWNT-fluoroelastomer nanocomposites measured by magnetic force microscopy (Reproduced from Ref. [117] with permisssion from Wiley Interscience)
matrix is paramagnetic. So when an MFM tip moves across the nanocomposite sample, many interactions, which include magnetic interactions, electrostatic interactions, Van der Waals interactions, short-range and capillary forces, take place. Another reason for the poor quality of MFM images of nanotube-polymer composites has been attributed to the basic working principle of MFM. Our group reported the preparation and properties of MWNT reinforced TPE gels [119]. Figure 13a and b display a pair of TEM images collected from NCTPE (nanocomposite thermoplastic elastomer) gels modified with 5 wt% MWCNTs taken from different sites on the same ultra-thin sample. The images of NCTPE gels exhibit morphology composed of a micellar of SEBS in hydrocarbon oil and MWCNTs. The irregularly shaped, long, and opaque features in both figures identify the incorporation of MWCNTs that measure on the order of 10–16 nm across and appear to flocculate into large-scale aggregates within the swollen midblock of the SEBS triblock copolymer network. In some areas as indicated by arrow, it is clear that these nanoparticles are flocculated into large-scale aggregates due to high solvent content and hinder to a lesser extent the bridging efficacy of individual copolymer molecules. The relatively high magnification image shows the more detailed examination of the NCTPE gel morphology and confirmed that the matrix consists of SEBS micelles measuring ca. 20 nm in core diameter uniformly dispersed throughout the hydrocarbon oil. It is clear from these related
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Fig. 13 TEM images of 5 wt% MWNT-based NCTPE gels at a low magnification and b high magnifications
images that the micelles do not exhibit any discernible indication of long-range order, a face- or, more likely, body-centered cubic lattice. The irregularly shaped dark features in both figures identify the incorporation of MWCNTs, which appear to exist as aggregates that measure on the order of 10–16 nm across. It is clear that MWCNTs are flocculated into large-scale aggregates due to high solvent content and hinder to a lesser extent the bridging efficacy of individual copolymer molecules, which possibly is dictated by poor dispersion.
4.2 Mechanical Properties The variation in tensile strength and modulus with increasing nanofiller loadings is shown in Fig. 14 (nanographite-reinforced fluoroelastomers), Fig. 15 (VGCNFreinforced chlorobutyl vulcanizates), Fig. 16 (MWNT-reinforced fluoroelastomers), and Fig. 17 (MWNT-reinforced TPE gels). From the figures it is observed that irrespective of the polymer matrix, there is a steady increase in tensile strength and modulus with increasing concentrations of the nanofillers. Tensile strength can be regarded as catastrophic tearing of cracks initiated by accidental flaws, microvoids, dewetting, or cavitation from the filler surface. If the elastomeric network is capable of dissipating the input energy into heat, then less elastic energy will be available to break this polymer network. Incorporation of fillers is the major source of energy dissipation. Increasing amounts of filler lead to a large
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Fig. 14 Variation in tensile strength and modulus with increasing nanographite concentrations in fluoroelastomer composites
Fig. 15 Variation in tensile strength and modulus with increasing carbon nanofiber concentrations in chlorobutyl vulcanizates
number of polymer chains adhering to the polymer surface, thereby leading to a greater probability of molecular slippage and thus increasing the fracture energy. From the figures it is also observed that the intensity of increase in tensile strength
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Fig. 16 Variation in modulus and tensile strength with increasing MWNT concentrations in fluoroelastomer composites
Fig. 17 Mechanical properties of NCTPE gels with MWNT: a tensile strength and b elongation at break
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Fig. 18 a High-resolution SEM microphotograph of exfoliated nanographite sheets; and b highresolution SEM microphotograph of nanographite sheets extracted from the polymer compound after solvent (MEK) extraction) (Reproduced from Ref. [116] with permisssion from Wiley Interscience)
is more pronounced when compared to modulus. The nanofillers used in the present study have been modified by acid treatment. This acid treatment of the nanofillers imparts functional groups like –COOH and –C=O on the surface, thereby increasing their oxygen functionality. The presence of these oxygen groups increases the polymer-filler interactions due the formation of complex physico-chemical bonds between the filler surface and the polymer matrix. Increased filler loading leads to an increase in polymer-filler interactions, thereby making a portion of the polymer matrix attach to the filler surface (the so-called bound rubber phenomenon). SEM microphotographs of raw nanographite powder and solvent extracted graphite from the polymer compound using methyl ethyl ketone (MEK) taken at the same length scales are shown in Fig. 18a and b, respectively. In Fig. 18b, adsorption of polymer chains onto the nanographite surface can be observed. The mechanical properties of TPE gels with increasing MWNT concentration are presented in Fig. 17a and b. For comparison, the tensile strength of triblock copolymer gels is 0.067 MPa and % elongation at break is 82, which corresponds to zero content in the figure. With a small addition of MWCNTs in TPE gels, there is a modest increase in tensile strength as the concentration increases. This can be explained by the fact that the incorporation of fillers is a major source of energy dissipation and due to the aspect ratio, thereby increasing the tensile strength of the TPE gels. As the particle size decreases, large surface areas become available.
4.3 Fractography The observation of fracture surface can give new insights regarding the failure mechanism of the composites. Although there have been many fracture surface studies in traditional filler (carbon black, carbon silica dual-phase filler, and silica) reinforced vulcanizates, seldom has it been investigated in nanofiller-reinforced
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elastomers. Figure 19a–d show representative scanning electron micrographs of the fracture surface of nanographite-reinforced fluoroelastomer composites at decreasing length scales. The crack growth phenomenon can be considered as a succession of elementary sequences from initiation to final failure. As the crack tip is stretched, cavities are induced by de-cohesion (or debonding) between fillers and the rubber matrix, forming microvoids. When these microvoids reach a critical volume, they coalesce into microcracks. With further application of stress, these microcracks propagate perpendicularly to the crack growth direction and eventually lead to tearing of the samples. Figure 19a reveals the presence of weld lines. The presence of weld lines can be attributed to parallel orientation of the graphite platelets in the plane of the weld line, thereby inhibiting the interdiffusion of polymer chains [122]. SEM micrographs at lower length scales (Fig. 19c, d) show the presence of aggregated and de-bonded nano-graphite particulates, which confirms the theory that the failure occurring during deformation of the composites is primarily via the coalescence of the voids forming critical cracks. From Fig. 19b it is also observed that the crack surface is rough, which indicates that crack propagation is difficult.
Fig. 19 Fracture morphology at different length scales in 3.5 phr nanographite-reinforced fluoroelastomer composites
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Fig. 20 Tensile fracture morphology of MWNT-reinforced fluoroelastomer nanocomposites (Reproduced from Ref. [117] with permisssion from Wiley Interscience)
Figure 20a and b show representative SEM microphotographs of fractured surfaces at two different length scales. From Fig. 20a the characteristic fracture morphology of nanofiller-reinforced elastomers, the so-called ‘‘cross-hatch pattern,’’ can be observed. This morphology is composed of numerous webs and steps of different sizes in particulate-reinforced vulcanizates. But in our case of MWNTreinforced fluoroelastomers, the numerous webs and steps appear to be nearly the same size. This can be attributed to the mechanism of crack propagation in MWNT-reinforced elastomers. Recent studies using H-NMR imaging demonstrated the presence of several microvoids in the elastomer matrix. These small voids tend to nucleate at the critical stage for crack growth, leading to lower mechanical properties. Stress–strain forces are amplified near a void, thereby creating nucleation sites for craze and crack growth. In a typical tensile test the rubber composite undergoes an immediate elastic deformation, after which the position of the crack changes slowly and follows a path of least resistance to its propagation. Increasing filler concentration in the polymer matrix reduces this propagation.
4.4 Dielectric Relaxation One of the main advantages of using exfoliated graphite is that it imparts good electrical conduction and dielectric properties to the composites. Special attention has focused on these properties. One of the most valuable tools for characterizing the relaxation behaviour of polymer systems is dielectric relaxation spectroscopy (DRS). DRS is a useful complement to the more customary mechanical methods of probing the viscoelastic properties of polymers [123]. Dielectric spectra reflect the same chain motions as the mechanical modulus; however, there is reduced interference due to symmetry from shorter time processes, thus making it more accurate than traditional dynamic mechanical analysis [124].
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4.4.1 Impedance Analysis Figure 21 shows the variation of the imaginary component of the complex impedance (Z00 ) with frequency as a function of increasing filler loading in exfoliated nanographite-reinforced fluoroelastomer vulcanizates. From the figure it is observed that with an increase in frequency, there is a gradual increase in complex impedance for all the vulcanizates, reaching a maximum in the region of 10 Hz at low filler loadings (0.5 and 1.5 phr) and around 100 Hz at higher loadings (2.5 and 3.5 phr). This can be attributed to the secondary relaxation of the polymer chains of the fluoroelastomer matrix. The observation of this additional peak in the imaginary component of impedance can also be attributed to the relaxation of the interfacial region between the polymer and the filler. Numerous examples of additional damping peaks have been cited in the field of particulate reinforced polymeric systems [125, 126]. This can be explained on the basis of the mechanical and viscoelastic properties of cross-linked and reinforced multiphase polymeric materials. Generally, polymer composites are cross-linked multiphase materials, the relaxation of which depends on molecular relaxation processes and morphology. Although these relaxations can usually be associated with each component, their appearance depends on the chemical and physical interactions of the two phases (i.e., filler and polymer matrix). The filler used in this study (exfoliated nanographite, prepared by acid intercalation technique), due to its high surface activity and the presence of oxygen, shows good interaction with the polymer matrix, thereby leading to the formation of a strong interphase. Fig. 21 Variation in complex impedance with increasing nanographite loadings (Reproduced from Ref. [118] with permisssion from Wiley Interscience)
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Fig. 22 Variation in real part of impedance with increasing nanographite loadings (Reproduced from Ref. [118] with permisssion from Wiley Interscience)
The thickness of this interphase is inversely proportional to the interfacial tension between the polymeric phases. A substantial portion of the polymer chains is expected to be immobilized on the filler surface, thus leading to the formation of regions of spatial heterogeneity. A polymer layer having a higher stiffness than the bulk polymer in the vicinity of the dispersed phase surface is created from restricted molecular mobility due to interactions between the phases [127]. So, as the filler loading increases, more and more polymer is adsorbed onto the filler surface, thereby making its relaxation difficult. A similar explanation can be extended to explain the decrease in the real part of impedance with frequency (Fig. 22). This trend of a continuous drop in impedance with applied frequency is characteristic of a pure capacitor.
4.4.2 Nyquist Plots Figure 23 shows the Nyquist plot [the relationship between the imaginary part of impedance (Z00 ) and the real part of impedance (Z0 )] of fluoroelastomer vulcanizates as a function of filler loading. It can be observed that increasing filler in the composite has a sizable effect on the dielectric properties of this system at all frequencies. From the figure it can also be observed that, irrespective of the filler loadings, the plots yield good semicircles, indicating the occurrence of polarization with a single relaxation time taking place, that is, a local mode process dominated. However, at higher filler loadings (2.5 and 3.5 phr), the semicircles did not reach the origin and had a small positive intercept on the Z0 axis, indicating a build-up of ions at the interphase between the filler and polymer matrix [128]. Several attempts have been made to interpret the impedance spectroscopy of polymerfiller systems using the resistance–capacitance parallel (R–C) circuit
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Fig. 23 Nyquist plot of fluoroelastomer-nanographite nanocomposites (Reproduced from Ref. [116] with permisssion from Wiley Interscience)
model. With increasing filler loading, the distance between the aggregates decreases. This gap can be approximated by a parallel plate capacitor with an area A, separation distance (d), and capacitance (C) eA/x, where e is the dielectric constant of the polymer. Each filler aggregate has a resistance Ra, the resistance within the aggregate. The impedance in a composite can be written as: Z ¼ Ras þ
Rcs xR2cs Cs j 2 2 2 1 þ x Rcs Cs 1 þ x2 R2cs þ Cs2
ð1Þ
The respective imaginary and real parts of impedance can be expresses as: Z 0 ¼ Ras þ
Rcs 1 þ x2 R2cs Cs2
and Z 00 ¼
xR2cs Cs 1 þ x2 R2cs Cs2
ð2Þ
and the dielectric loss tangent can be expressed as: tan d ¼
Z 00 xR2cs Cs ¼ 0 Z Ras þ Rcs þ x2 R2cs Ras Cs2
ð3Þ
Role of Different Nanoparticles Table 3 Radius and center for Nyquist plots in nanographite-reinforced fluoroelastomer vulcanizates
35 Filler loading (phr)
Center (x, 0)
0.5 1.5 2.5 3.5
3.424 2.273 4.358 4.462
9 9 9 9
105, 105, 104, 104,
Radius 0 0 0 0
From above equations, the relationship between Z0 and Z00 is: 2 2Ras þ Rcs 2 2 Rcs 0 Z þZ2 ¼ 2 2
1.98 0.87 4.42 4.34
9 9 9 9
105 105 104 104
ð4Þ
Therefore, a plot of Z00 and Z0 will give a half circle that has its center at {(2Ras ? Rcs)/2, 0} and a radius of Rcs/2. Wang et al. [129] proposed that because the circular curve of Z00 versus Z0 occurs only for the parallel resistor circuit, then the above analysis can be used to confirm the existence of the capacitor effect. The capacitor effect also confirms that the gaps between the nanographite control electron conduction via non-Ohmic contacts between the filler aggregates. The variation in the values of radius and center of the half-circle can also be used as a measure of the gaps between the filler aggregates. Using the above equations, the radius and center have been calculated and are tabulated in Table 3. It can be observed that with increasing filler loadings, the radius reduces and the center shifts to lower values, which implies increasing capacitance.
4.4.3 Percolation According to classical percolation theory, the percolation limit of any property of the material MP in a polymer reinforced with a filler and the volume fraction f of particles in the mixture above percolation threshold f c can be expressed as: MP ¼ MP0 ðf fc Þt
ð5Þ
where t is the critical exponent that is related to the dimensionality of the system and MP0 is the material property in pure polymer. Changing Eq. 5 to log scale, one can calculate the value of critical exponent t. A system obeying the percolation model should give a straight line: log MP ¼ t logðf fc Þ þ log MP0
ð6Þ
Two material properties—conductivity as obtained from dielectric relaxation spectroscopy (at various frequencies) and modulus values obtained from stress– strain measurements—have been used to study the percolation phenomenon and are plotted in Figs. 24 and 25, respectively. From Figs. 24 and 25 it can be observed that percolation is occurring at 2.5 phr in electrical conductivity (at all tested frequencies), above 1.5 phr in storage
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Fig. 24 Variation in AC conductivity with increasing nanographite loadings measured at increasing frequencies (Reproduced from Ref. [116] with permisssion from Wiley Interscience)
Fig. 25 Variation in modulus at increasing nanographite loadings (Reproduced from Ref. [118] with permisssion from Wiley Interscience)
modulus (at all temperatures tested), and above 2.5 phr in modulus as obtained from stress–strain curves. The values of the critical exponent t are shown in Table 4. From Table 4 it is observed that the value of the critical exponent t decreases with both increasing frequency and increasing temperature (as in storage modulus). A similar observation of a decrease in the value of t with increasing temperature was reported by Levya et al. [130], who attributed this phenomenon to the
Role of Different Nanoparticles Table 4 Value of critical exponent t
37 Property
Test conditions
Exponent t
Conductivity from DRS
100 Hz 1,000 Hz 10,000 Hz 100,000 Hz +22C
2.67 2.54 2.36 2.17 3.03
Modulus from stress–strain
occurrence of multiple percolation thresholds. A similar observation in carbonblack-reinforced polymer blends was also reported by Levon et al. [131] and theoretically proven by Roberts et al. [132].
4.5 Dynamic Mechanical Properties One of the main usages of butyl and halobutyl elastomers is as inner tubes in tires and vibration-isolating applications in the automotive industry. Therefore, in the present study, special attention focused on the dynamic mechanical properties of VGCNF-reinforced chlorobutyl vulcanizates.
4.5.1 Effect of Temperature on Loss Tangent Figure 26 shows the loss tangent spectra of chlorobutyl vulcanizates reinforced with VGCNF as a function of temperature. Increasing the amount of filler in the vulcanizates has no significant effect on the temperature or location of the maximum value of loss tangent (tandmax), but the magnitude of the peak decreases increasing filler loading. All samples showed the glass transition in the range of -29 to -34C. When a polymer is cooled through the glass transition region, the physical properties of the polymer in the non-equilibrium state (at temperature less than Tg), such as volume and enthalpy, gradually recover to new equilibrium values through configurational rearrangement of the polymer segments. The rate of the rearrangement or relaxation process depends on the local environment surrounding the relaxation entities and hence reflects the extent of environmental restriction on those entities. However, the magnitude of the Tg shift is marginal in our case. This can be explained on the basis of the ‘‘sluggish’’ nature of polyisobutylene relaxation dynamics. The glass-to-liquid transition primarily involves correlated local motions of a few backbone chains [124]. However, the segmental relaxation in polyisobutylene- (PIB-) based elastomers is weak, due to the steric hindrance by the methyl side groups. It has been suggested that low barriers to internal rotation about the skeletal bonds are a contributing factor in maintaining the segmental mobility necessary for elastomeric behavior. In polyisobutylene (PIB) (–C–(CH3)2–CH2–), the methyl groups bonded to alternate chain carbon atoms
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Fig. 26 Variation in loss tangent as a function of temperature in VGCNF-reinforced chlorobutyl vulcanizates
produce steric crowding. Although partially relieved by distortion of the (–C– CH2–C–) and (–CH2–C–CH2–) bond angles (127 and 109, respectively), it is still significant and hence PIB is not very flexible. Because butyl and halobutyl elastomers are copolymers of 97 wt% isobutylene and 3% isoprene, the relaxation in CIIR (Chlorobutyl Rubber) is predominantly dominated by the polyisobutylene chains. Further, the distribution of the end-to-end distance of chain segments between topologically adjacent cross-links plays a crucial role in the relaxation process [133]. Because the isobutylene content of CIIR has no chemically reactive site (unsaturated bond or double bond), and the only available reaction sites are Cl (1.35%) and the unsaturation of isoprene, the average chain length between the two adjacent reactive sites is very large. Consequently, the vulcanization network in CIIR is expected to be more homogeneous than those in other rubbers; thereby, the glass transition solely depends on polyisobutylene chain dynamics. Beyond the glass transition temperature (i.e., in the rubbery region) at higher VGCNF loadings (15 and 10 phr), it is observed that there is a small hump in the loss tangent curves. This can be attributed to the relaxation of the interfacial region of the polymerfiller in the composites. A similar behaviour is also observed in carbon-black-filled bromobutyl, chlorobutyl, and styrene butadiene rubbers where a similar phenomenon has been observed [134, 135].
4.5.2 Effect of Temperature on Storage and Loss Modulus The variation of storage modulus with temperature for elastomeric composites is a very important design criterion. Lewis and Nielsen [125] have pointed out that an elastomer composition with dynamic mechanical properties less sensitive to temperature changes is required for vibration isolation applications. Figure 27 shows the variation in storage modulus (E0 ) for different vulcanizates over a
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Fig. 27 Variation in storage modulus as a function of temperature in VGCNF-reinforced chlorobutyl vulcanizates
temperature range of -80 to +100C. In the glassy region (-80 to –30C), the augmentation in storage modulus E0 in the composite with the addition of fillers can be attributed to the hydrodynamic effect of the filler particle embedded in the polymer continuum. In the rubbery region (-30 to 100C), the polymer-filler and filler-filler networking (i.e., structure) and aggregate interactions have a pronounced effect on the value of E0 . Carbonaceous fillers such as carbon blacks and carbon silica dual-phase fillers affect the elastic properties of the vulcanizates more than the hydrodynamic reinforcement. The filler used in the present study, VGCNF, is also carbon-based filler that has been oxidized by treating with acid. This acid treatment increases the oxygen functionality of the fillers and has a profound effect on polymer-filler interactions, as widely reported by Serizawa et al. [136] in isobutylenebased elastomers. Figure 28 represents the temperature dependence of the loss modulus (E00 ) for various compositions. For all filled compositions, a distinct transition peak (a transition) is observed around -50C that can be attributed to conformational transitions occurring in the polyisobutylene backbone caused by micro-Brownian motions. The amount of filler has no significant effect on a peak location and a intensity. The amount of filler does not affect the TE00 max temperature (temperature at a peak); however, the intensity of a peak increases slightly with filler loading. Higher filler loadings result in a percolated network of filler particles that can influence relaxation on different scales. The percolation effect is usually considered effective for relaxation of longer time scales, such as the terminal relaxation observed in some rheological measurements. Recent dynamic mechanical experiments for composite solids seem to indicate that restriction effects do, in fact, result from the formation of a percolation network [137]. Nevertheless, reports regarding the restriction effects of percolated networks on the segmental relaxation are not fully conclusive, and the issue awaits more detailed and systematic study.
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Fig. 28 Variation in loss modulus as a function of temperature in VGCNF-reinforced chlorobutyl vulcanizates
4.5.3 Strain-Dependent Dynamic Mechanical Properties Although there have been many reports about strain-dependent dynamic properties (Payne effect) in vulcanizates reinforced with carbon-based fillers such as carbon blacks and carbon silica dual-phase fillers, it has never been reported in VGCNFreinforced elastomers. Because a typical elastomeric composite must withstand a wide range of temperatures in many practical applications, the Payne effect was studied at three temperatures in this study: -30C (near the glass transition temperature), room temperature (+30C), and high temperature (+70C).
Effect of Strain on tand and Storage Modulus at -30C The variation in tand and storage modulus with strain at -30C is shown in Fig. 29a and b, respectively. The storage modulus decreases and increases with increasing strain for all the compounds. All the vulcanizates show highest storage modulus (E0 ) at the lower strains. At the lowest strain of 0.07%, the threedimensional filler-filler and filler-polymer structure acts as a rigid unit against the imposed strain and hence gives the highest modulus value. The strain input associated at low strains is not sufficient to cause any significant change in network structure. The reinforcement at a moderate strain (\2%) is greatly affected by a disruption of the continuous network of filler that interpenetrates the rubber matrix. Under application of strain, the molecules of smaller chain lengths between the densely packed network points get oriented and form crystallites, whereas the molecules of much longer chain lengths would remain in random coil states. CIIR seldom exhibits strain-induced crystallization at normal operating temperatures, but recent studies have shown that butyl elastomers are capable of forming oriented lamellar structures at low temperatures [138]. They form a relatively high oriented amorphous fraction (*50%). The strain-induced
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Fig. 29 Effect of dynamic strain amplitude on a loss tangent and b storage modulus in VGCNF-reinforced chlorobutyl vulcanizates at 30C at increasing filler loadings
crystallites are in the form of a microfibrillar structure containing extended chain conformation. These crystallites are connected by oriented amorphous tie chains, which form a new network structure that is immobilized in a large volume of random amorphous chains. Luch et al. showed the disappearance of lamellar crystallites with increasing cross-link density, whereas recently Shimizu et al. showed that in uncrosslinked natural rubber, there is a growth of lamellar crystallites [139, 140]. It must be mentioned that the direction of orientation of these crystallites is controversial. Keller theoretically showed that the strain induced crystallization in polymers grow in the direction parallel to the chain axis, but Conradt showed the development of lamellar crystallites in the direction perpendicular to the molecular chain axis [141, 142]. However, the general consensus is that the strain-induced crystallization increases the storage modulus. Moreover, the
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filler particles are anisotropic in nature and get oriented in the direction of applied strain, thus leading to higher storage modulus values. Effect of Strain on tan d and Storage Modulus at Room Temperature (+30C) The variation in tand and storage modulus with strain at room temperature (+30C) is shown in Fig. 30a and b, respectively. The familiar phenomenon of a decrease in storage modulus and an increase in tand with increasing strains for all the compounds is observed. With increasing strain, it is observed that there is a decrease in the storage modulus. This can be explained on the basis of polymer-filler interactions, the desorption and reabsorption of the hard rubber shell surrounding the filler aggregate, or breaking and reforming of the effective cross-link in the vulcanizates-forming transition zone between the bound rubber and the bulk rubber. Fig. 30 Effect of dynamic strain amplitude on a loss tangent and b storage modulus in VGCNF-reinforced chlorobutyl vulcanizates at 30C at increasing filler loadings
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At lower strains, the three-dimensional filler-filler and filler-polymer structure acts as a rigid unit against the imposed strain and hence will give a higher modulus. The strain input associated with low strains is not sufficient to cause any significant change in network structure. However, at higher strains, this secondary structure tends to break down. Some researchers tend to equate ‘‘microstructure’’ with ‘‘agglomeration’’ or filler networking, and therefore filler agglomeration is also widely cited for the nonlinear behavior of filled polymer melts. Many researchers showed that the chemical structure of the polymer chains and fillers have a profound influence on relaxation behavior. Many studies have shown that the surface oxidation of carbon-based fillers leads to increased interactions with the polymer matrix [143]. At higher filler loadings, there may be two competitive phenomena occurring. Higher filler loading will give rise to the availability of a larger surface area of the fillers, thereby increasing the polymer-filler interactions. On the other hand, the secondary aggregate structures also increase at higher loadings, especially at 15 phr. Because the inter-particle potential is a strong function of the separation distance, this induces a steep dependence of the energy stored in the network of the particles on the macroscopic strain.
Effect of Strain on tan d and Storage Modulus at +70C Figure 31a and b show the variation of tand and storage modulus with strain at +70C. There is a decrease in storage modulus with increasing strain. But at all filler loadings, the decrease is much more pronounced when compared to -30C and room temperature (+30C). Increased filler loading in the composites form secondary aggregate structures. This may lead to new damping mechanisms such as particle–particle friction, particle-polymer friction where there is essentially no adhesion at the interface, and excess damping at the interface due to induced thermal stresses or changes in polymer conformation or morphology. Some recent studies show that the linear viscoelastic behavior of filled elastomers is governed by the glass transition gradient near the particle surface, combined with the arrangement of filler particles within the elastomermatrix [144]. If an increasing strain is applied to a reinforced system, its dynamic modulus decreases while the modulus of the nonreinforced matrix remains constant. Thus, the strain-softening of the glassy shell surrounding the solid particles is responsible for the Payne effect.
4.6 Rheological Properties One of the most important properties of TPE gels is the MOT (maximum operating temperature), which is characterized by the transition from a solid-like state to a viscoelastic liquid state. A rheological technique has been utilized in this study to find the MOT of MWNT-reinforced TPE gels. Hybrid gels were subjected to dynamic mechanical studies to discern the effect of nanoparticles on the
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Fig. 31 Effect of dynamic strain amplitude on a loss tangent and b storage modulus in VGCNF-reinforced chlorobutyl vulcanizates at 70C at increasing filler loadings
rheological behavior of hybrid triblock copolymer gels using a strain-controlled Rheometrics Mechanical Spectrometer (RMS 800, Rheometric Scientific, USA.), operated with 25-mm parallel plate geometry and a 1.5- or 2.5-mm gap height in the temperature range between 30 and 140C. The oscillatory shear measurements focused on the variations of elastic (in-phase) G0 the out-of-phase, G00 , as a function of temperature and frequency.
4.6.1 Effect of Frequency Measurements of the real and imaginary parts, G0 and G00 , of the shear modulus were also made as a function of frequency, w, of a small deformation strain of 1%
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for nanoparticles. The effect of frequency of NCTPE gels containing different amounts of MWCNTs is shown in Fig. 32. The results show that G0 exceeds G00 over the entire experimental frequency range, and G0 depends on the frequency at this 1% strain, which is typical of a physical gel. Here, we mean physical gels as a liquid-rich system exhibiting solid-like behavior, which the characteristics of showing flat mechanical spectrum in an oscillatory shear experiment [145]. The slight increase in G0 as the frequency increases may still be considered negligible. It can also be seen from Fig. 32 that the modulus increases as the concentration increases. The effect of frequency at higher temperature was also studied. The onset of maximum operating temperature, such as 70C for MWCNTs and 90C for parent TPE gels, was chosen to investigate the effect of temperature as a function of frequency (Fig. 33a, b). A further increase in temperature causes the moduli to become frequency dependent with increasing w. This only implies that the gel is no longer physical gel at these temperatures and the system loses its elasticity. In the case of NCTPE gels with MWCNTs, G00 is lower than G0 at lower frequency and then crossover in G0 and G00 is also observed as frequency increases, thereby indicating that the gels become viscous liquids. Comparable behavior was observed in TPE gels with 20% SEBS concentration, suggesting that just below this temperature, the gel network reinforcing TPE gels undergo some supramolecular rearrangement, perhaps similar to an orderdisorder transition (ODT). Note that the loss of gel behavior (i.e., the onset of frequency dependence in G0 ) is most pronounced at low frequencies, as is typically observed for physical gels. Fig. 32 Storage modulus G0 (closed symbols) and loss modulus G00 (open symbols) of NCTPE gels with MWCNTs presented as a function of oscillatory frequency (w) at a strain amplitude of 1% at 30C
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Fig. 33 Variation in storage modulus G0 (closed symbols) and G00 (open symbols) as a function of oscillating frequency at 70C for (a) 5 wt% MWNT and (b) 20 wt% MWNT loaded TPE gels
4.6.2 Effect of Temperature Figure 34 shows the change in G0 with temperature heated throughout the melting range for NCTPE gels with 5 wt% MWCNTs and parent TPE gels (Fig. 35). The parent TPE gel consists of four characteristics region: (1) an initial plateau over which G0 remains relatively constant or slightly increases with increasing temperature, which means that the rubbery PS domains become glassy; (2) an abrupt reduction in G0 attributed to a lattice disordering transition in the vicinity of 75C; (3) a small second plateau in G0 at 100C; and (4) a steep decrease in G0 , signifying Fig. 34 Dynamic storage modulus as a function of temperature at strain amplitude of 1% for parent TPE gels (squares) and NCTPE gels with 5 wt% MWCNT (circles)
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Fig. 35 Dynamic storage modulus G0 (closed symbols) of NCTPE gels with MWCNTs of 0.5 wt% (hexagon), 1 wt% (triangle), 3 wt% (down triangle), and 5 wt% (star) and dynamic loss modulus (G00 ) (open symbols)
the collapse of the copolymer network at ca. 108C. The existence of a second plateau reflects the persistence of a residual, load-bearing nanostructure that either remains intact or develops upon increasing the temperature. Montensen et al. [146] have demonstrated that similar TPE gels exhibit a high-temperature, bodycentered cubic phase that possesses an unusually high degree of nanostructured order. The nanosystem TPE gels have only three major characteristic regions without the second plateau observed in the parent TPE gels, that is, without any morphological change during the heating procedure. The precipitous decrease in G0 at very low temperature was observed for all NCTPE gels, which are much lower than their parent TPE gels. An abrupt reduction in G0 was reduced in the vicinity of 75C. This reduction in G0 indicates the collapse of the triblock copolymer network in which the lattice disordering transition. If the MWCNT loading increases, G0 also increases.
4.6.3 Effect of MWNT Concentration on the Storage Modulus At a temperature between 30 and 40C below the gel point, the NCTPE gels have the property of elastic moduli, where G0 [ G0 . This phenomenon was found at every nanoparticle composition. This indicates that at ambient temperature, a physical network is still present despite the addition of nanoparticles. In this regard, the values of the plateau G0 extracted from the data between 30 and 40C is displayed in Fig. 35 as a function of nanoparticle. The parent TPE gel with 20 wt% SEBS concentration was included here as reference material for comparison.
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Low concentrations yield low G0 values with their parent TPE gels. MWCNTs are probably flocculated into large-scale aggregates due to high solvent content and hinder to a lesser extent the bridging efficacy of individual copolymer molecules, which are possibly dictated by dispersion problems encountered, and are further explained in morphological observation. The further addition of 3–5 wt% concentrations to SEBS/oil gel likewise promotes a modest increase in G0 , which may be due to the high aspect ratio nature of this material.
4.7 Thermal Property Thermo-gravimetry (TG) is a suitable method to evaluate the thermal degradation properties of elastomers. The thermal stability of TPE has been extensively studied because of the great importance of this group of materials. Reinforcing the inorganic filler improves the thermal stability of thermoplastic elastomers and plastics. Because the reinforcing of nanoparticles improves several properties, which promises wide applicability, it would be interesting to study the thermal stability of NCTPE gels. The derivative thermogravimetric (DTG) and thermo-gravimetric (TG) curves for NCTPE gels with MWCNTs and parent TPE gels are shown in Fig. 36a and b. The TG curve has two distinct regions of weight loss that are reflected in two peaks (maxima) in the DTG curve, implying that at least two stages of degradation occur in this sample. The initial degradation in stage I results primarily from the decomposition of the swollen EB (ethylene butylene) block in hydrocarbon oil, while stage II is due to the SEBS matrix. For each sample, the
Fig. 36 a DTG and b TGA of NCTPE gels with increasing amounts of MWNTs in the parent TPEG
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thermogram revealed that the DTG plot (Fig. 36a) shows a maximum rate of weight loss (Fig. 36b). It can be clearly observed from these figures the thermogram shift toward higher temperature as the heating rate increases. The shift of temperature is more pronounced at lower temperatures. It appears that the particles reside in the region of the EB block swollen by a high content of oil, as seen in TEM image, thereby increasing the distinct region of oil degradation temperature. It can also explained that may be due to the filler effect that is prominent and retards the degradation of TPE gels and the retardation effect is attributed to the interaction between nanoparticles and macroradicals generated during the degradation process. However, the second peak that appears at higher temperature remains constant with addition of MWCNTs. The weight loss for swollen EB containing oil and SEBS matrix remains constant with the addition of MWCNTs. The residual yields of the NCTPE gels increased with increasing MWCNT content, indicating that thermal decomposition of the polymer matrix was retarded in the NCTPE gels with higher residual yield. This result may be attributed to a physical barrier effect due to the fact that MWCNTs would prevent the transport of decomposition products in the polymer nanocomposites. Similar observations have been reported that the thermal stability of polypropylene/layered silicate nanocomposites was improved by a physical barrier effect; hence it is enhanced by ablative reassembling of the silicate layer [147]. Comparing the residue of TPE gels without MWCNTs, there is little residue because the component of the gels consists only of carbon and hydrogen.
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In Situ Synthesis of Rubber Nanocomposites Massimo Messori
Abstract The preparation and characterization of rubber based nanocomposites prepared by in situ generation of inorganic oxides by means of the hydrolytic sol– gel process are reviewed in the present chapter. The sol–gel approach has been applied to several rubber matrices to prepare reinforced vulcanized and unvulcanized rubbers. Several synthetic procedures are presented while the most investigated filler is silica obtained by hydrolysis and condensation of tetraethoxysilane. The effects of the different preparation conditions and of the filler content are generally discussed in terms of morphology (investigated by electron microscopy and small angle X-ray scattering) and mechanical properties (modulus, strength and extensibility). The mechanical properties of the in situ filled nanocomposites are generally better than those of the corresponding materials prepared with the conventional mechanical mixing of preformed particulates and elastomers. This enhancement is generally attributed to a lower tendency to filler–filler aggregation due to a lower particle surface interaction resulting from the ‘bottomup approach’ of the sol–gel process applied to the preparation of organic–inorganic hybrid materials.
1 Introduction The improvement of the mechanical properties (reinforcement) of elastomeric materials by addition of rigid fillers represent one of the most important aspects in the field of rubber science and technology [1]. The concurrent enhancement of M. Messori (&) Dipartimento di Ingegneria dei Materiali e dell’Ambiente, Università di Modena e Reggio Emilia, Via Vignolese 905/A, 41125 Modena, Italy e-mail:
[email protected] V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_2, Ó Springer-Verlag Berlin Heidelberg 2011
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stiffness (elastic modulus) and, possibly, of elongation at break due to the presence of rigid particles derives from hydrodynamic effects mainly depending on the filler volume fraction but also affected by the filler shape factor (aspect ratio). The interactions between rubber and filler, which can be increased if good dispersion and distribution of the particulate are achieved, play a fundamental role for the increase of elastic modulus. The surface characteristics of the particles (presence of reactive/functional groups, wettability and surface energy) and the chemical nature of the rubber represent the key parameters for these interactions. Concerning the field of fillers having dimensions in the nanometric scale, polymer matrix nanocomposites have attracted extraordinary attention in the last decade on the basis of their excellent mechanical and barrier properties compared to the conventional microcomposites, usually at very low filler content. Layered silicates, ceramic nanoparticles (such as silica, titania, zirconia, etc.), carbon nanofibers and nanotubes are typical examples of materials used as nanosize reinforcing additive [2]. The usual method for the preparation of nanocomposites is based on the top-down approach according to which preformed nano-objects are dispersed within the polymeric matrix by physical–mechanical dispersion and distribution (melt or solution mixing). Even if very attractive from an industrial point of view, this method presents some severe limitations related to the difficulties to obtain an effective dispersion due to the strong tendency to particles aggregation phenomena and the significant increase of melt viscosity because of the complex rheology of nanocomposite systems. In particular, in the specific case of rubber composites, it is well known that carbon black represents the most effective reinforcing additive for elastomeric compounds notwithstanding the strong limitation due to the black coloration imparted to the final part. Alternatively to carbon black, also silica is another filler widely employed in the rubber industry, thanks to some advantages deriving from its use such as high tear strength, good abrasion resistance and reduction in heat build-up [3]. The main drawbacks associated with the use of conventional silica as reinforcing additive instead of carbon black are the higher compound viscosity, the incompatibility of silica and rubber, the more difficult mixing and processing, the longer vulcanization time, the lower cross-linking density and the absorption of curative additives by silica. Most of these disadvantages are due to the very strong interaction occurring among silica particles and caused by the hydrogen bonding of the silanol groups present onto the silica surface. This interaction inhibits an uniform dispersion of the filler within the rubber matrix giving rise to the formation of particles’ aggregates. Silica agglomeration and silica–rubber incompatibility are generally reduced by using different types of silanes as coupling agents. In order to minimize these problems and alternatively to the usually employed processing technique (mechanical mixing) for the preparation of filled vulcanized rubber, the in situ generation of silica (or other inorganic oxides) by using the sol– gel process has been recognized as a novel and interesting technique for the preparation of rubber composites.
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The classical aqueous (or non-hydrolytic) sol–gel process [4] consists of a two-step hydrolysis-condensation reaction starting with metal alkoxides M(OR)x, typically tetraethoxysilane (TEOS), according to the following scheme: Step 1: Hydrolysis
MðORÞx þ n H2 O ! MðORÞxn ðOHÞn þ n ROH Step 2: Condensation
MOH þ HOM ! MOM þ H2 O ðwater condensationÞ and/or MOR þ HOM ! MOM þ ROH ðalcohol condensationÞ The presence in the reactive system of an organic oligomer or polymer (bearing or not suitable groups reactive towards to the sol–gel process) leads to the formation of organic–inorganic hybrid structures composed of metal oxide (silica or other) and organic phases intimately mixed each other. This synthetic procedure belongs to the so-called ‘bottom-up’ approach for the preparation of hybrid materials and, depending on the experimental conditions, permits the synthesis of composite structures in which the dimensions of the dispersed phase are under 100 nm (nanocomposites). The optical, physical and mechanical properties of these nanocomposites are strongly dependent not only on the individual properties of each component, but also on important aspects of the chemistry involved such as uniformity, phase continuity, domain size and the molecular mixing at the phase boundaries. The morphologies of the hybrid materials are strictly dependent on the characteristics of the organic polymer such as the molecular weight, the presence and the number of reactive functionalities as well as the solubility of the polymer in the sol–gel system. The nature of the interface between organic and inorganic phases is generally used to grossly divide these materials into two distinct classes. In Class I, organic and inorganic components are simply embedded and only weak bonds (hydrogen, dipolar and van der Waals bonds) give the cohesion to the whole structure. On the contrary, in Class II materials, the two phases are linked together through strong chemical bonds (covalent or ionic bonds). Several applications of organic–inorganic hybrid nanocomposites prepared by sol–gel and other procedure have been already extensively reviewed [5–8]. It is also important to underline that the use of the sol–gel procedure for the preparation of rubber matrix nanocomposites could present some peculiarity mainly due to the possibility to vulcanize a preformed organic polymer (‘green’ rubber).
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The majority of the reported papers deals with the preparation and characterization of elastomers filled with in situ generated silica obtained from TEOS as precursors while a much more limited literature on other inorganic oxides such as titania, zirconia and alumina is available. Several synthetic procedures are proposed and the characterization is mainly devoted to mechanical and/or thermomechanical properties, such as quasi-static tensile measurements or dynamic mechanical analysis, and to morphology investigation by means of scanning and transmission electron microscopy and/or small-angle X-ray scattering. One of the main targets is the development of nanocomposites with improved properties with respect to similar materials obtained with the conventional mechanical mixing of rubber and preformed nanoparticles. In the present chapter a general overview on the in situ generation of silica or other reinforcing fillers within elastomeric phases reported in literature is presented and discussed for different elastomeric matrices.
2 General Synthetic Strategies From a generic point of view, the in situ generation of an inorganic oxide phase (usually having a spherical or quasi-spherical shape) by means of the sol–gel process within an elastomeric polymer can be carried out according to different strategies, as step-by-step schematically described in the following. Solution procedure (procedure A) (i) dissolution of both the metal alkoxide (metal oxide precursor) and unvulcanized rubber in a common solvent; (ii) addition of water (directly added or absorbed from the external humid atmosphere), sol–gel catalysts (basic or acidic) and vulcanization ingredients (optional) and activation of the sol–gel process at a given temperature and for a given reaction time; (iii) removal of solvent and by-products (usually water and alcohols) by evaporation or precipitation in a non-solvent; (iv) vulcanization of the filled ‘green’ rubber (optional). Swelling of unvulcanized rubber (procedure B) (i) immersion and swelling of films or sheets of unvulcanized rubber in metal alkoxide; (ii) immersion of the swollen rubber in a basic or acidic aqueous solution and activation of the sol–gel process at a given temperature and for a given reaction time; (iii) mechanical mixing (usually by using a conventional two-roll mill) of the filled ‘green’ rubber with the vulcanization ingredients and vulcanization (optional).
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Swelling of vulcanized rubber (procedure C) (i) immersion and swelling of films or sheets of vulcanized rubber in metal alkoxide; (ii) immersion of the swollen vulcanized rubber in a basic or acidic aqueous solution and activation of the sol–gel process at a given temperature and for a given reaction time. It can be anticipated that the procedure A (in solution) ensures a highly homogeneous dispersion of the in situ generated filler within the rubber matrix also for high reinforcing metal oxide contents. The swelling of unvulcanized rubber (procedure B) is highly attractive especially for industrial applications thanks to its solvent-free approach while the swelling of vulcanized rubber (procedure C) is very simple and can be in principle applied to already moulded three-dimensional articles (even if limitations in the maximum amount of filler generated can be present due to the cross-linked structure of the swollen rubber).
3 Rubber Based Nanocomposites 3.1 Polydimethylsiloxane Based Nanocomposites Polydimethylsiloxane (PDMS) is the most important member of the class of polysiloxanes and presents some interesting properties thanks to the presence of Si–O–Si linkages in the polymeric backbone such as good thermal stability, water repellency, excellent resistance to oxygen, ozone and UV-light, anti-stickiness and low chemical reactivity. As drawbacks, PDMS is characterized by very poor mechanical properties, in particular tensile strength, and it requires the addition of reinforcing additives (usually mineral fillers) in order to achieve the mechanical properties required depending on the final application. PDMS is traditionally reinforced with silica and the filler–matrix interactions are ensured through hydrogen bonding between silanol groups of silica surface and oxygen atoms of PDMS macromolecules. Surface modification of silica particles permits the optimization of the interactions between mineral filler and polymer. Due to the concurrent presence of Si–O groups in both silica and PDMS, it is not surprising that these represent one of the first systems investigated in the field of in situ generation of silica within polymeric matrix, as indicated by the pioneeristic works of Mark [9, 10].
3.1.1 PDMS-Silica Nanocomposites According to a slight modification of the above described procedure C (swelling of vulcanized rubber), Rajan et al. [11] reported a method (named controlled hydrolysis, CH) for the in situ generation of silica particles in which, after swelling
62 Table 1 PDMS/SiO2 composites: sol–gel catalysts, filler content and mechanical properties (adapted from Ref. [11])
M. Messori Catalyst
Silica Ultimate Elongation (wt%) strength (MPa) (relative length) at rupture
None Dibutyltin diacetate
0.0 5.4 8.9 14.2 6.0 8.0
Dibutyltin dilaurate
0.082 0.538 0.818 0.712 0.543 0.831
1.77 1.68 1.63 1.09 2.01 1.93
of cross-linked PDMS in TEOS, the required water of hydrolysis was simply absorbed from the air and the catalyst was generated from tin salts (dibutyltin diacetate or dilaurate). These PDMS nanocomposites had a silica content up to 14 wt% and were characterized by an unusually high transparency (if compared to similar materials obtained with aqueous ammonia as catalysts) as quantitatively judged by UV–VIS spectroscopy. Electron microscopy showed that silica domains were very small (30–50 nm in diameter) and well dispersed, as expected from the transparency of the composites. Tensile stress–strain measurements indicated that the particles provide very good reinforcement. Data reported in Table 1 indicate that the ultimate tensile strength and Young’s moduli increased with higher silica content and elongation at break remained almost the same of the unfilled PDMS (up to a silica content of about 9 wt%). The dependence of silica particle dimensions on PDMS cross-linking density, silica content and catalyst concentration was investigated by small-angle X-ray scattering [12]. Of particular interest were the relationships between particle size and molecular weight of the network chains (Mc, mesh sizes), amount of filler introduced and catalyst concentration. Silica particle sizes were smallest for the smallest values of Mc, presumably due to constraining effects deriving from the very short network chains. At fixed Mc and filler concentrations, higher catalyst concentrations gave larger particles. Increase in silica concentration generally had little effect on particle size at low and high loadings, but markedly increased sizes at intermediate levels (10–20 wt%), presumably caused by coalescence of isolated small particles into considerably larger aggregates. Films comprised of PDMS and in situ generated SiO2 were applied on silicon, aluminum and polystyrene substrates. The surfaces in contact with air and with a substrate were investigated by using several surface technique [13]. The hybrid sample surfaces generated in contact with air were characterized by a silica-free PDMS top layer of about 2 nm; in the surfaces in contact with the substrates SiO2 was located at or just beneath the outermost atomic layer. In contact with polar liquids such as water, polar hydroxy groups present at the surface of SiO2 can easily stretch out to the outer-most atomic layer. Quite surprisingly, no significant correlation was found between the roughness of the surfaces and the amount of in situ generated SiO2 present in the materials.
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Fig. 1 TEM image of PDMS/SiO2 composites (silica content 10 phr) (reproduced with permission from Ref. [14])
Bokobza et al. [14, 15] applied the same synthetic procedure to obtain similar composites by starting from a cross-linked vinyl-terminated PDMS. Authors found that the in situ generated silica structures were uniformly dispersed within the elastomeric matrix (see Fig. 1) but different morphologies were obtained as a function of the used tin-based catalyst and thus of the type of growth processes. The hydrophilic character of the silica surface was responsible of extensive interaction with PDMS chains leading to a significant improvement in the mechanical properties of the composites. An interesting comparison among PDMS reinforced with different types of fillers has been published by Bokobza [16]. Three different techniques of incorporation of silica were compared: the conventional mechanical mixing process, in situ filling process and introduction of spherical colloidal silicas (Stöber silicas). In addition, results obtained with other types of fillers (layered silicates and fibrous clays) were reported and discussed. Among the different types of reinforcement of PDMS imparted by silica particles, the in situ filling process was by far the most efficient. As reported in Table 2, the modulus, ultimate stress and extensibility increased by increasing the filler loading. Moreover, larger increases were observed over the silica particles percolation threshold. Table 2 PDMS/SiO2 composites: mechanical properties (adapted from ref. [16]) Compound Stress at 100% Stress at 200% Tensile strength Elongation strain (MPa) strain (MPa) (MPa) at break (%) Unfilled sample 10 phr SiO2 18 phr SiO2 30 phr SiO2 40 phr SiO2
n.r. 1.44 2.49 5.19 7.99
n.r. 4.98 7.86 10.35 16.38
0.43 7.64 10.70 21.54 27.34
91 242 240 313 275
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3.1.2 PDMS-Titania and PDMS-Zirconia Nanocomposites The same CH method was used by Murugesan and Mark [17, 18] for the preparation of PDMS containing in situ generated titania and zirconia starting from different alkoxides and the obtained materials were compared to those prepared according the so-called conventional ‘water excess process’ (WE, based on the direct addition of stoichiometric water to the cross-linked PDMS swollen in the corresponding metal alkoxide). The composites prepared by the CH method were characterized by a greater amount of filler (higher conversion of alkoxides to zirconia or titania) for the same reaction time. The optical transparency of the prepared materials is reported in Fig. 2 and UV–VIS analysis showed that the transmittance values of PDMS/ZrO2 were significantly higher than those of PDMS/TiO2 (47 and 28% for composites containing 6 wt% of filler and 16 and 0.01% for composites containing 28 wt% of filler, respectively) according to the fact that titania particles formed small aggregates compared to zirconia, as evidenced by small-angle X-ray scattering analysis and leading to a higher opacity. Thermogravimetric analysis showed that PDMS/ZrO2 composites had a lower thermal stability with respect to PDMS/TiO2 composites, which were inherently more stable due to the formation of different phases at high temperatures. Concerning the mechanical properties, all the PDMS/ZrO2 and PDMS/TiO2 composites (independently on the preparation method) had mechanical properties much improved with respect to the unfilled elastomer. The reinforcement of zirconia filled composites was greater compared to titania filled composites even for similar amount of filler presumably due to the smaller dimension of ZrO2 particles, consistently with the above discussed transparency.
Fig. 2 Transparency comparison of PDMS/ZrO2 and of PDMS/TiO2 composites (reproduced with permission from Ref. [18])
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3.1.3 PDMS-(mixed oxides) Nanocomposites Wen and Mark [19] reported the in situ generation of silica–titania mixed oxide within a PDMS network (previously cross-linked with TEOS). TEOS and tetrabutyl titanate (TBT) were the precursors of the correspondent oxides. Since the TEOS hydrolysis is extremely slow compared to titanium alkoxides, hydrolysis of a mixture of them generally results in phase segregation leading to a physical mixture of two oxides instead of a true mixed oxide. This problem has been avoided carrying out a partial hydrolysis of pure TEOS with subsequent addition of the more reactive TBT. The sol–gel process has been activated by placing the swollen PDMS in 2 wt% diethyleneamine aqueous solution at room temperature leading to composites with a silica–titania content ranging from 10 to 22 wt%. TEM analysis carried out on PDMS/SiO2–TiO2 composites showed a filler average diameter of 20–25 nm, with a relatively narrow diameter distribution and with very little aggregation of particles. A further interesting observation was the increase of the particle size by decreasing the cross-linking density of the rubber, as already evidenced for PDMS/SiO2 nanocomposites [12]. The analysis of equilibrium swelling data on the basis of the Kraus’ theory [20] suggested that the composites were of the adhering type in which the rubber matrix is restricted by the filler through attachments onto the filler surface. Concerning mechanical properties, stress–strain isotherms represented as plots of modulus against reciprocal elongation according to the Mooney-Rivlin equation [21, 22] indicated that, compared to the silica-filled PDMS networks, the mixedoxide based composites had better extendibility and had upturns which occurred at higher elongations (but were smaller because of the lower reinforcing effect provided by TiO2). The presence of in situ generated silica–titania mixed oxides also increased the onset temperature for thermal degradation of PDMS, on the contrary of that showed by physical mixing of preformed silica and/or titania nanoparticles. The same authors reported a further study [23] on the preparation and characterization of PDMS/SiO2 and PDMS/TiO2–SiO2 composites having inorganic contents ranging from 30 to 100 wt% and in which the sol–gel reaction (conversion of TEOS in SiO2 and of TBT in TiO2 and corresponding mixed oxides) occurred simultaneously with the cross-linking of PDMS due to the co-condensation of hydroxyl-terminated PDMS and TEOS. Experimental data indicated that this last process was dominant and that the majority of the PDMS was incorporated in the silica network. The addition of PDMS was found to shorten the gelation times and to increase the rates of increase in modulus of the network structure. The impact strength and the fracture surfaces of the materials were studied with the finding that the presence of PDMS significantly increases their impact strength and ductility, as shown in Fig. 3. The growth processes and resulting structures of the reinforcing fillers were investigated by small-angle X-ray scattering by Breiner and Mark [24]. The systems were found to yield dense particles with fractally rough surfaces and the results were used to interpret mechanical properties of these composites.
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Fig. 3 Dependence of impact strengths on PDMS content for PDMS/SiO2 composites. The impact strength were obtained from a the Charpy pendulum impact test and b the falling-weight impact test (reproduced with permission from Ref. [23])
The reinforcement with mixed oxides seems to synergistically combine the best properties of titania and silica: the titania provides additional reinforcement at low strains and allows for higher elongations to be obtained, whereas the silica contributes to the reinforcement over the entire range of allowable strain and gives rise to the very desirable upturn in modulus at high elongations. Wen and Mark [25] also reported the precipitation of silica–zirconia and silica– alumina mixed oxides (at a concentration ranging from 10 to 22 wt%) into PDMS networks using a sol–gel approach according the above described method. The resulting filled networks were found to have very good mechanical properties. In comparison with networks filled only with silica, these materials had good extensibilities as well as high strengths. Filler particle diameters were generally several hundred angstroms, but also in this case were found to decrease with increase in cross-linking density of the networks. The distributions of particle size were relatively narrow, and there was very little particles aggregation. The presence of in situ generated silica–titania and silica–zirconia mixed oxides also improved the thermal stability of the PDMS.
3.2 Natural Rubber Based Nanocomposites Bokobza and co-workers [26, 27] reported the preparation and the characterization of natural rubber (NR) composites containing in situ generated silica particles by hydrolysis and condensation of TEOS before the vulcanization of the rubber matrix in solution (according to the procedure A) and after the vulcanization of the rubber (according to the procedure C) with or without the addition of a silane coupling agent (bis(3-triethoxysilylpropyl) tetrasulfide, TESPT). The silica content of the prepared composites varied from 8 to 21 phr. Tensile stress–strain experiments showed that, at a given strain, the stress increased by increasing the silica content showing a significant reinforcing effect due to the presence of in situ generated silica.
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It is generally accepted that the increase of the stiffness (modulus) deriving from the generation or incorporation of active filler within a polymeric matrix is related to two factors: (i) a hydrodynamic effect arising from the presence of rigid particles and (ii) an increase in the cross-linking density due to polymer–filler interactions. The first term can be quantitatively described by the following Guth and Gold equation [28, 29]: G ¼ G0 ð1 þ 2:5 / þ 14:1 /2 Þ where G and G0 represent the moduli of the composite and the unfilled polymeric matrix, respectively, and / represents the volume fraction of the filler. As shown in Fig. 4, the stress–strain curve calculated by using the Guth and Gold equation was lower with respect to the experimental curve indicating that, besides a hydrodynamic reinforcement, the presence of filler–rubber interactions must be considered. Similar conclusions were also reported analysing the stress– strain data according to the Mooney-Rivlin plots and on the basis of equilibrium swelling tests. Authors also demonstrated that silica particles generated before rubber vulcanization (according to the procedure A) inhibit the cross-linking reaction of the rubber compound by sulphur. The presence of a filler networking structure, due to the aggregation of silica particles via the silanol groups present on the surface, was suggested by the Mooney-Rivlin plot analysis. On the other hand, the tendency to form a particle–particle network was avoided when silica was generated in the already vulcanized rubber (according to the procedure C) and small and well dispersed particles were observed. Pissis and co-workers [30, 31] investigated the molecular dynamic of the above discussed materials. Broadband dielectric relaxation spectroscopy investigations reported in these papers demonstrated that, in addition to the a relaxation related to the glass transition of the rubber matrix, a slower a relaxation was observed and assigned to polymer chains close to the polymer/silica interphase whose mobility
Fig. 4 Experimental and calculated (Guth and Gold equation) stress–strain curves for NR/SiO2 composites (reproduced with permission from Ref. [26])
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M. Messori
is restricted due to interaction with the silica surface. Authors reported an estimation of the thickness of this interphase of about 2.1–2.4 nm. Poompradub and co-workers [32, 33] prepared in situ silica reinforced NR by means of the procedure in solution (procedure A) with n-hexylamine as basic catalysts without any subsequent vulcanization and investigated the preparation conditions and the rheological and mechanical properties of the prepared rubber composites (having a silica content ranging from 10 to 70 phr). Different solvents were used for the synthesis (CCl4, CHCl3 and tetrahydrofuran, THF) and it was found that the lower amount of silica in rubber matrix was generated when CCl4 and CHCl3 were used compared with THF. The best behaviour of THF was attributed to its higher water solubility and polarity. As expected, they observed an increase in Mooney viscosity by increasing the in situ generated silica content. Interestingly, at the same filler content the viscosity of composites containing in situ generated silica was lower compared to that of rubber reinforced with preformed silica dispersed by mechanical mixing. This behaviour has been attributed to the presence of a lower amount of silanol groups onto the surface of in situ generated silica with respect to commercial preformed silica, as further supported by dynamic mechanical analysis which suggested that the silica–silica interactions of the in situ generated silica were weaker, resulting in a better dispersion in the rubber matrix. Also the moduli and compression set values of NR filled with in situ generated silica were improved with respect to the commercial silica ones. The reinforcement has been modelled by using the Guth and Gold equation with a shape factor f = 2.53 (according to the morphology evidenced by TEM analysis). Authors concluded that the volume fraction of filler, the particle–particle interaction and the anisotropy of silica particles represent the three main factors affecting the reinforcing effect due to the in situ generation of silica within the NR matrix. Ikeda and Poompradub [34–37] reported the preparation and the characterization of NR composites containing in situ generated silica particles from TEOS in the presence of amines as catalyst before the vulcanization of the rubber (thus according to the procedure B). As a general procedure, sheets of NR having a thickness of about 1 mm were prepared by two-roll milling and immersed in TEOS at room temperature or slightly higher and for different times. The swollen NR sheets were further immersed in aqueous solution of amine-based catalyst at 40°C for 75 h in order to carry out the conversion of TEOS to silica. NR sheets containing the in situ generated silica were mixed with sulphur-based vulcanizing agents in a two-roll mill at room temperature. Then, the rubber compound was moulded by hot-pressing to produce vulcanized sheets of about 1 mm of thickness. As expected, the amount of in situ generated silica was found strongly dependent on the amount of TEOS in the swollen NR and on the type of catalyst. Authors found that the polarity and the basicity of amine catalyst were crucial for controlling the conversion of TEOS to silica. In this respect, primary alkylamines with suitable hydrocarbon segments (n-hexylamine, n-heptylamine and n-octylamine) produced the highest amount of in situ generated silica (up to about 80 phr). Due to its higher solubility in water, n-hexylamine was found to be the
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Fig. 5 Stress-elongation curves for NR/SiO2 composites (filled circle in situ generated SiO2; filled square SiO2 mechanically mixed) (reproduced with permission from Ref. [34])
most effective catalyst. Almost independently on the type of catalyst, all the silica particles were generated in the level of nanometric size (average particle size ranging from 25 to 45 nm) and with spherical shape. The tensile properties of the materials have been investigated and compared to those of rubber unfilled and filled with preformed nanosized silica by using the conventional mechanical mixing method. Stress–strain curves reported in Fig. 5 indicate that for the same filler content (71 phr) and for comparable network chain densities, the in situ silica reinforced rubber showed the lower stress at the elongation up to about 200% and higher stress at the elongation beyond about 200%. The high modulus at low elongation for conventional silica filled rubber has been assumed to be due to the formation of larger aggregates of particles suggesting a better dispersion in the case of in situ silica reinforced rubber. On the other hand, the high modulus at high elongation for in situ silica filled rubber has been attributed to a stronger interaction between NR matrix and in situ generated silica particles. The reinforcement effect due to in situ generated silica was compared to that of carbon black (both low structure and high structure carbon black stocks) [38]. Physical and mechanical properties of NR reinforced with in situ generated silica were in between those of carbon black filled composites and conventional silica filled composites (both prepared by mechanical mixing). Ikeda and co-workers [39, 40] also reported a detailed investigation on the effect of a coupling agent (c-mercaptopropyltrimethoxysilane, c-MPS) on NR filled with in situ generated silica. The addition of c-MPS was found to improve the reinforcement effect. The diameter of in situ generated silica particles was in
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M. Messori
the range 20–45 nm and their dispersion within the rubber matrix became homogeneous by adding the silane coupling agent thanks to a reduction of silica– silica interactions. The concurrent use of in situ generated silica and c-MPS also prevented the delay of sulphur curing and increased the wettability of NR onto the particles’ surface, which resulted in the increase of reinforcement of the vulcanized composite. Ikeda et al. [41, 42] recently reported an interesting investigation on the applicability of 3D-TEM, a technique that combines transmission electron microscopy with computerized tomography, to visualize and analyze the threedimensional state of NR containing in situ generated silica. Tangpasunthal and co-workers [43, 44] described the in situ generation of silica within NR latex. Mixtures of TEOS and other alkyltriethoxysilanes (TESPT, vinyltriethoxysilane, ethyltriethoxysilane and i-butyltriethoxysilane) were directly added to a commercial grade NR latex, containing ammonia and thus permitting the direct formation of silica through the sol–gel reaction thanks to the presence of a basic pH. Compounds with sulfur-based vulcanizing agents were subsequently prepared by two-roll milling and vulcanized by hot pressing. The conversion of alkoxysilanes to silica within the rubber was almost complete for TEOS but decreased when the alkyl group of the alkyltriethoxysilane increased in size. Silica particles with sizes between 100 and 500 nm and evenly dispersed without extensive aggregation were obtained. The presence of TESPT, a coupling agent widely used in rubber industry, resulted in an increase of the mechanical properties and the rate of sulphur cure. Hardness, tensile and tear properties of vulcanized NR reinforced with in situ generated silica were higher than those of similar composites prepared by conventional mechanical mixing. Among the different alkyltriethoxysilane investigated, vinyltriethoxysilane seemed to be the most promising taking into account the high enhancement in tensile modulus and resistance to tear of the vulcanized rubber.
3.3 Epoxidized Natural Rubber Based Nanocomposites Epoxidized natural rubber (ENR)/silica nanocomposites were prepared by Bhowmick and co-workers [45, 46] according to the solution procedure A. TEOS (as precursor for the in situ generation of silica) and ENR were dissolved in THF and the sol–gel process was activated at room temperature under hydrochloric acid catalysis. Alternatively to NR, ENR with adequately high epoxy content was chosen as rubber matrix taking into account that epoxy groups, randomly distributed along the chain backbones, give higher glass transition temperature and, most importantly, increased polarity and thus stronger interaction with the in situ generated silica. In fact, it is well known that under acidic conditions, the epoxy groups are likely to open up as a diol moiety which can undergo intermolecular hydrogen bonding with the silanol groups. After
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solvent elimination, some samples were vulcanized with benzoyl peroxide or dicumyl peroxide. The composite films appeared transparent up to 50 wt% of TEOS loading. Dispersion of the discrete silica particles having dimensions ranging from 15 to 100 nm (by increasing the TEOS content) was observed. Infrared spectroscopic studies indicated the occurrence of chemical interaction within the rubber–silica interfaces which was further supported by the insolubility of the respective samples in THF under the ambient conditions. Dynamic mechanical analysis further corroborated the interactive nature of the in situ generated silica with the ENR matrix. All the ENR/SiO2 composites showed a significant improvement in mechanical properties with increased TEOS loading within the rubber. A maximum increment in tensile strength (200%) and tensile modulus (170%) with respect to the unfilled ENR was observed for the uncured composites with the highest TEOS content (50 wt%). Further reinforcement was noticed when the rubber in the nanocomposites was cured with either benzoyl peroxide or dicumyl peroxide. The dicumyl peroxide cured hybrid composites displayed 112% improvement in tensile strength over the control cross-linked rubber sample, probably due to synergisms of nanosilica reinforcement and cross-linking of the rubber phase in the hybrids. The effect of polymer-silica interaction on THF swelling and dynamic mechanical properties have been investigated and compared to those of acrylic rubber/silica and poly(vinyl alcohol)/silica composites [47]. A further comparative study on structure–property relationship of hybrids prepared under different pH levels has been published by the same authors [48]. The silica particles were formed in the nanometer scale (average diameter \100 nm) at low pH (equal or lower than 2) beyond which aggregation occurred, although the conversion of TEOS to silica was not strictly influenced by the various pH conditions. These nanocomposites were optically clear and showed superior mechanical reinforcement over the microcomposites containing aggregated silica structures with lower optical clarity. Furthermore, the nanocomposites exhibited higher storage modulus both at the glassy and the rubbery regions as compared to those of microcomposites. As evidenced by dynamic mechanical analysis, the tand peak heights were also minimum and the glass transition temperature shifted to higher temperature for those nanocomposites. An interesting and alternative synthetic procedure has been proposed by Hashim et al. [49, 50] according to which ENR sheets were first cross-linked with 3-aminopropyltriethoxysilane (APS) by hot pressing. After vulcanization, ENR sheets were swollen in TEOS and subsequently subjected to a sol–gel reaction in butylamine aqueous solution. Silica content was up to about 30 wt% according to conversions of TEOS to silica higher than 60%. The obtained sol–gel vulcanized rubber were more rigid and stronger than a typical sulphur-cured vulcanized ENR containing comparable amount of silica. Comparative stress–strain and dynamic mechanical property analysis suggested that chemicals bonds were formed between the silica particles and the rubber network thanks to the dual reactivity of APS with respect rubber vulcanization and sol–gel reaction.
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3.4 Isoprene Rubber Based Nanocomposites Messori and Bignotti [51] recently published a work on the in situ generation of silica from TEOS within synthetic cis-1,4-polyisoprene (isoprene rubber, IR), the synthetic counterpart of NR. According to the solution procedure A, the sol–gel process was activated in solution of toluene, in the presence of all components and by absorbing the water required for the hydrolysis of TEOS from the external humid atmosphere. After solvent elimination, vulcanization of the rubber was attained for some sample by thermal activation of dicumyl peroxide as vulcanizing agent. The conversion of TEOS to silica was quantitative for samples with a low initial TEOS concentration while for initial concentrations higher than 20 wt% the yield in silica was in the range 60–80% almost independently on the formulation and the curing conditions (presence or absence of coupling agent, vulcanization or not of IR phase). The in situ generated silica particles were homogeneously dispersed in the vulcanized rubber with a spherical shape and an average dimension which increased from a few nanometers to the submicron scale (300–400 nm) by increasing the concentration of inorganic phase. Swelling experiments evidenced that good polymer-filler adhesion was observed in the presence of coupling agents, while no definite conclusions could be drawn for vulcanized materials produced in the absence of the coupling agent. The dynamic mechanical behaviour of the various elastomers became increasingly nonlinear for silica contents higher than 20 wt%. In addition, only in this range of compositions the filler exerted on the low amplitude storage modulus a remarkable reinforcement, which was related to the silica content through a power law with exponent a = 4, in agreement with the prediction of a model proposed by Huber and Vilgis [52] as reported in Fig. 6.
Fig. 6 Huber and Vilgis model: excess storage modulus at the lowest strain amplitude investigated (cmin) versus silica volume fraction (U0) for IR/SiO2 nanocomposites (open diamond unvulcanized rubbed; filled circle vulcanized rubber; open circle vulcanized rubber with coupling agent). The straight line represents the best-fitting power law with exponent a = 4 (see text) (reproduced with permission from Ref. [51])
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A synthetic procedure based on direct addition of water to the reactive medium and of octyltriethoxysilane as coupling/surfactant agent at different reaction times was recently proposed by Messori et al. [53]. For long enough reaction time, the sol–gel process led to a quantitative conversion of TEOS precursor to silica with the obtainment of a homogeneous distribution of spherically shaped particles. The delayed addition of octyltriethoxysilane was effective for controlling the average size and the aggregation phenomena of the in situ generated silica. Dynamic mechanical analysis carried out on filled IR showed a significant reinforcement (in terms of storage modulus increment) with respect to the pristine elastomer. On the other hand, both swelling and extraction tests suggested that the sol–gel process perturbed the vulcanization process of IR leading to a slight decrease of the crosslinking degree of the rubber matrix.
3.5 Styrene-Butadiene Rubber Based Nanocomposites Ikeda et al. [54, 55] described the preparation of styrene-butadiene rubber (SBR) reinforced with in situ generated silica from TEOS under different catalytic conditions. Composites were prepared by swelling of vulcanized SBR in TEOS or in TEOS-THF mixture in the presence of basic (n-butylamine) or acid (hydrochloric acid) catalysts for the sol–gel reaction. Hydrochloric acid was found to be inadequate for the sol–gel reaction in the system comprised of vulcanized SBR swollen in TEOS because the aqueous solution of hydrochloric acid was not well dissolved in TEOS and, as a consequence, silica was formed only in the surface layer of the vulcanized rubber [55]. In order to avoid this problem, pre-swelling of vulcanized SBR was carried out in THF before the sol–gel reaction and, under these conditions, in both the acidic and basic aqueous solutions mixed with THF, the in situ formation of silica occurred homogeneously in the SBR matrix. The size of in situ generated silica was observed to be influenced by the cross-linking density, i.e. the larger the cross-linking density, the smaller the size of in situ formed silica particles, according to what already reported for other systems [12, 19, 25]. Comparing with the silica-filled vulcanized SBR prepared by conventional mechanical mixing, the homogeneity of dispersion of the silica particles was found to be important for the reinforcement of vulcanized rubber. Stress–strain curves for different materials are reported in Fig. 7, from which a marked improvement of mechanical properties (modulus and tensile strength) is evident for composites prepared under basic condition with respect to SBR unfilled and filled with conventional silica. In addition, the size of silica particles obviously affects the reinforcement of the rubber. The larger the in situ silica particles, the better the mechanical properties found in this study. Interestingly, and contrary to what generally expected, the tand peaks (Tg) detected by dynamic mechanical analysis were found to decrease by about 2–4°C with respect to unfilled rubber. Authors proposed that this to be due to the swelling in TEOS, which may contribute to the disentanglement of the SBR chains in the
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Fig. 7 Stress–strain curves of vulcanized a SBR; b SBR/ SiO2 (mechanically mixed, SiO2 22 wt%); c SBR/SiO2 (pre-swelling in THF, acid catalysis, SiO2 13 wt%); d SBR/SiO2 (pre-swelling in THF, basic catalysis, SiO2 24 wt%) and e SBR/SiO2 (basic catalysis, SiO2 23 wt%) (reproduced with permission from Ref. [54])
vulcanizate followed by lowering of Tg. The plasticization of the rubber by residual oligomers from the sol–gel reaction might contribute to this lowering of Tg. Ikeda et al. [56] also reported an investigation on the effect of TESPT as coupling agent in the preparation of SBR filled with in situ generated silica. The presence of TESPT resulted in a much higher reinforcing efficiency with respect to conventional mechanical mixing and the in situ method without TESPT. The higher reinforcing efficiency was attributed to the formation of a silica–rubber network, which also changed the dynamic mechanical behaviour of the vulcanized rubber. Transmission electron microscopy analysis showed in situ silica incorporation of very fine particles in comparison to the sol–gel process without TESPT. Similarly to the above described approach used by Hashim for ENR [49, 50], de Luca and co-workers [57, 58] prepared a Class II hybrids comprising of epoxidized SBR, TEOS (as silica precursor) and APTS (as coupling agent) in order to enhance the interaction between organic and inorganic phases. The epoxy groups of the rubber reacted with the amino groups of APTS forming an intermediate amino-silane/rubber compound which in turn was reactive towards the sol–gel reaction of TEOS to silica. The pre-reaction between epoxidized SBR and APTS was carried out in THF and TEOS and water were subsequently added (the alkaline pH for the sol–gel synthesis was provided by the APTMS itself). Large amounts of silica were incorporated using combinations of the inorganic precursors TEOS and APTS. Non-solubility of the materials in THF indicated the formation of a network, the microstructure of which varied according to the concentrations of the inorganic precursors employed. The mechanical properties
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increased considerably by increasing the amount of silica incorporated and SEM analysis revealed the presence of phase separation for high TEOS content. In situ silica generation from TEOS has been also applied in the case of blends of SBR and reclaim rubber (RR) by De and co-workers [59]. They presented an interesting comparison of mechanical properties of nanocomposites comprised of SBR/RR blends containing silica both in situ generated and mechanically mixed. They also investigated the effect of the presence of TESPT as coupling agent. Authors concluded that the mechanical properties of conventional nanocomposites were better than those of in situ generated ones in the absence of TESPT but a reverse trend was observed in the presence of the coupling agent.
3.6 Acrylonitrile-Butadiene Rubber Based Nanocomposites In situ silica reinforcement was applied to acrylonitrile-butadiene rubber (NBR) vulcanizates which were swelled in TEOS and subsequently soaked in an aqueous solution of ethylenediamine [60]. The amount of in situ generated silica within the NBR vulcanizates (conversion of TEOS to silica of 63%) was limited due to the high polarity of NBR and the resulting low degree of swelling of NBR in TEOS. The presence of c-MPS in the rubber vulcanizate increased the conversion of TEOS to silica during the sol–gel reaction (conversion of TEOS to silica higher than 90%), compared to the system without c-MPS. The obtained silica particles were very fine and very homogeneously dispersed. The same authors also proposed the in situ generation of silica from TEOS within NBR vulcanizates pre-mixed with conventional silica both in the presence or not of c-MPS [61]. They observed that the reinforcement efficiency tended to increase with the increase of mechanically pre-mixed silica. Both transmission electron and scanning electron microscopies showed that the simultaneous use of pre-mixed silica and c-MPS promoted the formation of large silica particles and clusters with a relatively good dispersion by the sol–gel reaction of TEOS in the rubber vulcanizate as further evidenced by the results of hysteresis measurements. This behaviour has been attributed to the surface modification of conventional silica by the sol–gel reaction of TEOS and the presence of c-MPS which worked as a dispersion agent for silica particles.
3.7 Butadiene Rubber Based Nanocomposites In situ silica filling of butadiene rubber (BR) was carried out by the sol–gel process using TEOS [62]. BR was sulphur-cured and the resultant cross-linked BR was firstly swollen in TEOS and subsequently immersed in an aqueous solution of n-butylamine at 30°C for 24 h and at 50°C for 72 h to activate the sol–gel reaction of TEOS. The in situ generated silica was homogeneously dispersed in the rubbery
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Fig. 8 Stress–strain curves of vulcanized a BR; b BR/ SiO2 (mechanically mixed, SiO2 22 wt%); c BR/SiO2 (in situ generated, SiO2 16 wt%) (reproduced with permission from Ref. [62])
matrix with spherical shape and size in the range 15–35 nm. As already reported for other systems, the size of in situ generated silica was influenced by the crosslinking density, i.e. the larger the cross-linking density, the smaller the size of in situ silica formed in the vulcanized rubber. Concurrently, the interaction between filler and BR matrix seemed to become larger. Compared to the conventional silica-filled BR vulcanizate, which was prepared by mechanical mixing of the silica particles, the vulcanized rubber with the in situ generated silica showed better mechanical properties (in terms of both modulus and strength, as shown in Fig. 8). Mark and Zhou reported the preparation and characterization of trans1,4-polybutadiene (tPBD) reinforced with in situ generated silica after swelling of the cross-linked tPBD with TEOS in the presence of a tin-based catalyst and by absorbing the required water from the external humid atmosphere [63]. The silica dimensions varied from below 100 nm at low silica content (\2 wt%) to 2–5 lm at higher content (up to 20 wt%). Concerning the mechanical properties, small amounts of silica (up to 2 wt%) increased the extensibility and stress at rupture with a subsequent increment of toughness.
3.8 Acrylic Rubber Based Nanocomposites Following an approach very similar to those of ENR, Bhowmick and Bandyopadhyay [46, 64] reported the preparation and the characterization of acrylic rubber (ACR) reinforced with in situ generated silica by means the acid-catalyzed
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hydrolysis-condensation of TEOS in solution of THF (procedure A). The proportion of TEOS was from 0 to 50 wt% (macrophase separation occurred beyond this last maximum content of TEOS). For comparison, composites were also prepared with preformed silica up to 30 wt% of its loading. After solvent elimination, some sample was vulcanized with benzoyl peroxide or with a mixed curing system comprised of hexamethylenediamine and ammonium benzoate. All the hybrid composites prepared by sol–gel technique were transparent while, on the contrary, the composites of ACR with mechanically mixed preformed silica were opaque at all compositions. Infrared spectroscopic analysis revealed the absence of significant shifts in peak position for carbonyl absorption in all the hybrid composites, suggesting that a homogeneous dispersion of silica particles within the organic matrix occurred without any significant interaction between the two phases. This conclusion was further supported by the complete dissolution of the uncured hybrid in THF at room temperature. Morphological analysis revealed the presence of discrete spherical silica particles with an average diameter that increase from 20 to 90 nm by increasing the TEOS content (from 10 to 50 wt%). The storage modulus detected by dynamic mechanical analysis above glass transition temperature increased by increasing the TEOS content according to the reinforcing effect of nanosilica network structure present in the rubber matrix. The tand peak height was gradually reduced, the peak broadened and the Tg values of the composites shifted towards higher temperatures by increasing the inorganic filler content in the composites. Concerning mechanical properties of the hybrid composites, the tensile strength increased by increasing the silica content. For the same silica content, the tensile strength of composites filled with in situ generated silica was significantly higher than that of conventional composites. The nanolevel mixing of in situ generated silica particles and the formation of strong Si–O–Si network within the ACR matrix were responsible for greater reinforcement in the rubber matrix. In the case of conventional silica filled rubber, the bigger sized silica particles did not provide higher surface area for interactions such as that of in situ generated silica and hence they were not as effective as the former. On curing the rubber phase, with either benzoyl peroxide or hexamethylenediamine/ammonium benzoate, the tensile strength was further increased, although benzoyl peroxide cured samples showed lesser improvements than the corresponding mixed cross-linking system. In order to verify the effect of the polarity of the rubber matrix on the properties of this class of composites, Bhowmick and co-workers [65, 66] synthesized acrylic copolymers and terpolymers by bulk polymerization of ethyl acrylate (EA), butyl acrylate (BA) and acrylic acid (AA) and used them as rubber matrix for the preparation of silica filled composites by using the usual procedure based on the acid-catalyzed hydrolysis-condensation of TEOS in THF solution. Authors reported a morphological investigation showing the presence of silica particles with an average diameter lower than 100 nm. The average diameter of silica particles was lower in the case of high extent of polarity and hydrophilicity of the rubber matrix, suggesting that these molecular parameters control the in situ
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generated silica particles size. In all the cases, no significant chemical interaction occurred at the silica–rubber interface, as depicted from spectroscopic analysis (FT-IR and NMR). In general, the mechanical properties increased by increasing the polarity of the matrix due to better rubber-silica interaction and uniform nanosilica dispersion. Also in this case, cross-linking further improved the mechanical properties.
3.9 Ethylene–Propylene–Diene Monomer Rubber Based Nanocomposites Ethylene–propylene–diene monomer (EPDM) rubber was modified with TESPT in an internal mixer in order to graft triethoxysilyl groups onto the polymeric backbone [67]. Triethoxysilyl-grafted EPDM sheets were swollen with TEOS and subsequently immersed in n-butyl amine aqueous solution to in situ generate silica particles. The silica filled EPDM rubber was finally mixed with vulcanizing agents and cured by hot pressing. Authors concluded that TESPT fragments present in the macromolecular chain of EPDM acted as nucleation sites for the growing of silica which in turn led to the formation of a strong silica-EPDM network chemically bonded to the rubber matrix. The pendant triethoxysilyl groups on the EPDM backbone played a fundamental role in uniformly dispersing silica particles within the rubber. For similar silica loadings, EPDM modified with in situ generated silica showed superior reinforcing efficiency with respect to materials prepared by ex situ process (mechanical mixing of precipitated silica within the rubber).
3.10 Other Rubber Based Nanocomposites Matêjka and co-workers [68, 69] reported some studies on rubbery epoxy resins reinforced with in situ generated silica from TEOS and with a silica content ranging from 6 to 22 wt%. The epoxy resin and the hardener were diglycidyl ether of bisphenol A (DGEBA) and polyoxypropylene-diamine (JeffamineÒ D2000), respectively. Long flexible polyether chain lead to a rubbery cured material and moreover it solubilizes siloxane structures formed in the sol–gel process resulting in a transparent hybrid. Authors reported several preparation methods: (i) one-stage process, in which all reactants were mixed in iso-propyl alcohol solution and reacted simultaneously; (ii) two-stage ‘simultaneous’ process, in which TEOS was pre-hydrolyzed under acidic conditions in iso-propyl alcohol solution and subsequently added to DGEBA and JeffamineÒ D2000 to activate the simultaneous formation of both organic and inorganic networks;
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(iii) two-stage ‘sequential’ process with preformed epoxide network, in which the cured epoxy resin was firstly prepared by stoichiometric reaction between DGEBA and JeffamineÒ D2000 and subsequently swelled in the sol–gel solution (TEOS, water and acid in iso-propyl alcohol) to activate the in situ generation of silica. Compact and large silica aggregates (100–300 nm in diameter) were observed in the case of hybrids obtained with one-stage process, whereas open polymeric structures of smaller aggregates (50–100 nm in diameter) composed of fractal particles were formed during the two-stage procedure with the pre-hydrolyzed TEOS. Moreover, this two-stage process resulted in a much faster gelation of the silica. In these ‘simultaneous’ processes, the silica network was formed faster than the epoxide one and its structure was not influenced by the presence of the epoxide and amine. However, in the case of the two-stage ‘sequential’ process, the preformed organic network suppressed growth of the silica aggregates by inter-particle condensation and the relatively smallest silica domains (10–20 nm) were formed. Authors concluded that the synthetic procedure controls the reaction mechanism, final structure and morphology. In particular, contrary to what shown by basic and neutral catalysis, acid catalysis promotes rapid hydrolysis of siloxane groups which results in the high content of silanol groups and an extensive grafting to the epoxide network. Grafting between organic and inorganic phases resulted in the formation of an interphase epoxide layer with reduced mobility and increased glass transition temperature. Concerning the mechanical properties, an increase in modulus by two orders of magnitude was achieved at a low silica content (10 vol%). Interestingly, dynamic mechanical analysis revealed the presence of a co-continuous morphology of the epoxy matrix and of the silica phase (and the silica-glassy epoxide phase) continuously extending through the macroscopic sample. Sunada and co-workers [70] reported the synthesis of alkoxysilane-modified polychloroprene latex by the emulsion copolymerization of 2-(3-triethoxysilylpropyl)-1,3-butadiene and chloroprene. This latex was mixed with unmodified polychloroprene (CR) latex and TEOS to obtain composites by sol–gel reaction in the latex under basic catalysis. After vacuum drying, the composites were compounded with curing agents by two-roll milling and subjected to vulcanization by hot pressing. Electron microscopy showed that the silica particles in unvulcanized composites had various diameters ranging from 0.1 to 0.6 lm, and their size became larger with the decrease of the silica content. Compared to similar composites prepared by conventional mechanical mixing with preformed silica, vulcanized CR/SiO2 composites obtained with in situ generation of filler showed that the tensile modulus and tear strength improved with an increase of the amount of modified CR. The metallocene-based poly(ethylene-octene) (POE) elastomer, which was developed using a metallocene catalyst by Dow and Exxon, has received much attention because of its peculiar properties such as uniform distribution of comonomer content and narrow molecular weight distribution.
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Wu et al. [71, 72] reported a systematic investigation on composites comprised of POE or POE grafted with maleic anhydride or acrylic acid (POE-g-MAH or POE-g-AA, respectively) as elastomeric matrix and in situ generated silica or silica–titania mixed oxides. The composites were prepared by addition of solutions of (i) TEOS, water and hydrochloric acid or (ii) silicic acid, tetraisopropyl orthotitanate (TTIP), water and hydrochloric acid to POE or POE-g-MAH or POE-gAA melted in an internal mixer at a temperature of 160–170°C. FT-IR, 29Si-NMR and XRD analysis showed that Si–O–C, Ti–O–C, Ti–O–Ti, Si–O–Si and Si–O–Ti linkages were formed in silica and silica–titania reinforced elastomers supporting the expected formation of covalent bonds between organic and inorganic phases and of SiO2–TiO2 mixed oxides. It was found that there are maximum values of tensile strength and glass transition temperature at about 10 wt% of inorganic filler (see Table 3). Authors explained this behaviour assuming that an excess of SiO2 or SiO2–TiO2 particles might cause separation between the organic and inorganic phases with a reduction of their compatibility. From the reported data it is interesting to observe that superior properties were obtained in the case of hybrids based on POE grafted with maleic anhydride or acrylic acid (POE-g-MAH or POE-g-AA) with respect to pure POE. The higher polarity and the presence of reactive groups (anhydride or carboxylic) in the grafted elastomer allowed the formation of stronger chemical bonds between organic and inorganic phases. The literature concerning the in situ generation of inorganic oxides through the sol–gel process within thermoplastic elastomers (TPE) is relatively limited. Lai et al. [73, 74] reported the preparation of thermoplastic polyurethane (TPU) reinforced with in situ generated silica starting from TEOS as precursor or from polysilicic acid extracted from an aqueous sodium metasilicate solution with THF. A solution procedure (procedure A) with acidic catalysis (hydrochloric acid or acetic acid) was adopted in both approaches to prepare the hybrid materials. Spectroscopic and dynamic mechanical analysis showed the presence of interfacial interactions between the organic and inorganic phases. Storage modulus increased at all concentrations of silica due to its reinforcement effect while the tensile strength exhibited a maximum value for filler content of 10–15 wt%. In contrast, the cutting strength decreased, probably due to a reduction of the energy dissipation from silica as physical cross-links. The acetic acid catalyzed system
Table 3 Glass transition POE/SiO2 Filler content TS Tg (wt%) (°C) (MPa) 0 3 10 20
-60 -58 -55 -58
27 28 31 26
temperature and tensile strength (TS) values POE-based composites POE-g-MAH/SiO2 POE/SiO2–TiO2 POE-g-AA/SiO2–TiO2 Tg (°C)
TS (MPa)
Tg (°C)
TS (MPa)
Tg (°C)
TS (MPa)
-58 -56 -52 -56
15 38 52 27
-60 -58 -56 -58
27 28 32 28
-59 -54 -48 -57
17 37 56 30
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showed better optical properties than the hydrochloric acid catalyzed system. The modified process with extraction showed improved mechanical properties and better optical properties than the conventional sol–gel process without extraction.
4 Concluding Remarks The sol–gel process has been established as a versatile, effective and innovative method for the in situ preparation of rubber nanocomposites. The great majority of the published studies concerns the generation of silica from TEOS within the most common elastomers even if some example of in situ generation of other oxides such titania, zirconia, alumina or mixed oxides is present in the literature. Different synthetic strategies can be applied to obtain materials with tailored properties, in particular the in situ generation of inorganic fillers can be obtained in solution or directly in the rubber, either swollen in metal alkoxides or in the melt state. Depending on the preparation conditions, the vulcanization of the rubber matrix can be carried out before, simultaneously or after the sol–gel process and the resulting morphologies and ultimate properties are generally affected by the specific synthetic procedure. As a general observation, finer morphologies (lower filler dimensions) are realized by increasing the cross-linking density of the vulcanized rubber presumably due to constraining effects deriving from the organic network. It seems also very promising that the mechanical properties (modulus, strength and extensibility) of the in situ filled nanocomposites are generally better than those of the corresponding materials prepared with the conventional mechanical mixing of preformed particulates and elastomers. This enhancement is generally attributed to a lower tendency to filler–filler aggregation due to a lower particle surface interaction resulting from the ‘bottom-up approach’ of the sol–gel process applied to the preparation of organic–inorganic hybrid materials. Future developments on in situ synthesis of rubber nanocomposites should involve further systematic studies on the possibility to use silica precursors alternative to TEOS as well as the incorporation of metal oxides different than silica, exploring the different reinforcing effect of titania, zirconia or other. Also the use of non-hydrolytic sol–gel process, which is characterized by the absence of water as reactant for the generation of inorganic oxides, should be an interesting field of research and development.
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72. Wu, C.S.: Synthesis of polyethylene-octene elastomer/SiO2–TiO2 nanocomposites via in situ polymerization: properties and characterization of the hybrid. J. Polym. Sci. Polym. Chem. 43, 1690–1701 (2005) 73. Lai, S.M., Wang, C.K., Shen, H.F.: Properties and preparation of thermoplastic polyurethane/ silica hybrid using sol–gel process. J. Appl. Polym. Sci. 97, 1316–1325 (2005) 74. Lai, S.M., Liu, S.D.: Properties and preparation of thermoplastic polyurethane/silica hybrids using a modified sol–gel process. Polym. Eng. Sci. 47, 77–86 (2007)
Part II
Characterization & Properties
Relaxation Phenomena in Elastomeric Nanocomposites G. C. Psarras and K. G. Gatos
Abstract Elastomeric nanocomposites are technologically important engineering materials, mostly because of their thermomechanical and electrical behaviour. Relaxation phenomena arising in rubber nanocomposites include contributions from both the elastomeric matrix and the presence of nanofiller, and in many cases reflect the interactions between matrix and filler. Dynamic mechanical analysis and broadband dielectric spectroscopy are two mutual complementary experimental techniques, which record the relative response of the tested material under the influence of a harmonically time varying mechanical or electrical field. Occurring relaxations originate from molecular dynamic effects, interfacial phenomena and phase changes. Studying relaxation phenomena provide valuable information upon the structure–property relationships of the nanosystems. Elastomeric nanocomposites can be classified according the employed matrix to non-polar and polar. Additionally, can be classified according the aspect ratio of the used filler. Suitably selecting the type and the amount of nanoinclusions the service performance of nanocomposites can be tailored. Abbreviations ABS AMIC BDS
Antilock braking system 1-Allyl-3methyl imidazolium chloride Broadband dielectric spectroscopy
G. C. Psarras (&) Department of Materials Science, School of Natural Sciences, University of Patras, 26504 Patras, Greece e-mail:
[email protected] K. G. Gatos Megaplast S.A. Research & Development Center, 4 Makedonias Str., 16672, Athens, Greece e-mail:
[email protected] V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_3, Ó Springer-Verlag Berlin Heidelberg 2011
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BR CB CNT CP DEA DMA DSC ENR EPDM FGPNC FHT HNBR HXNBR IP LS MMT MQ(NMR) MWCNT MWS NMR NR OMMT phr POSS PUR SBR SDS TEM UDPNC VFT XNBR XRD
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Butadiene rubber Carbon black Carbon nanotube Cross polarization Dynamic electrical analysis Dynamic mechanical analysis Differential scanning calorimetry Epoxidized natural rubber Ethylene-propylene-diene rubber Functionally graded polymeric nanocomposite Sodium fluorohectorite Hydrogenated nitrile rubber Hydrogenated carboxylated acrylonitrile butadiene rubber Interfacial polarization Layered silicates Montmorillonite Multiple quantum (NMR) Multi-wall carbon nanotubes Maxwell–Wagner–Sillars Nuclear magnetic resonance Natural rubber Organically modified montmorillonite Parts per hundred rubber Polyhedral oligomeric silsesquioxanes Polyurethane Styrene butadiene rubber Sodium dodecyl sulfate Transmission electron microscopy Uniformly dispersed polymeric nanocomposite Vogel–Fulcher–Tamann Carboxylated nitrile rubber X-ray diffraction
1 Introduction The term relaxation phenomenon or effect refers to the process, which undergo a system or substance under the influence of an external field by changing its original equilibrium state to a new dynamic one. After removing the exciting force the system relaxes back to its initial equilibrium state [1]. The time scale of the system’s response function compared to the exerting time of the external field governs the relaxation process. If the time scale of the disturbing field is much
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longer than the system’s response time, then response function is allowed to reach its maximum value. If, on the other hand, the time scale of the field is short relative to the system’s response time, then limited variation of the equilibrium position is expected to occur, holding the response function to low or even minimum value. The application of an external electric field on a polar material constitutes a typical example of such a process. If the time of applying the field is long enough for the dipoles to be orientated to the direction of the field then the material becomes fully polarized and its response function (polarization) takes maximum value. On the contrary, if the time needed by the dipoles to become parallel to the applied field is much longer than the time of field application a weak disturbance of the initial state will take place and the achieved polarization of the material will remain low [1, 2]. Relaxation phenomena occur in many different materials (such as glasses, liquids, suspensions, disordered solids, ceramics, magnetic materials, polymers, composites, etc.). Among them, elastomers and their micro- and nano-composites have attracted the interest of the scientific community. Studying their relaxation phenomena provide valuable information concerning the relaxation time, the distribution or not of relaxation times, the activation energy of the process, the temperature and/or pressure dependency of the effect and the dynamics of molecular entities with respect to the applied stimulus and the exerted thermodynamic variables. Rubber technologists belong to the pioneers who utilized fillers of nanodimensions (e.g. carbon black) however, recent practice reveals their stiffness in accommodating the late advances in nanotechnology. The reason behind is likely the multi-component rubber vulcanization recipes, along with the complex chemistry involved during curing. Therefore, restrained utilization of novel fillers within the conservative and rather empirically built formulations took place. Since the initiation of the industrial applications of elastomers in the nineteenth century, fillers played a dominant role. Reinforcement of rubber stocks with chalk, talc, kaolin, zinc oxide and CB were among the first, which were exploited [3]. Usually, the fillers are classified as active and inactive. In the first case the fillers aim mainly to reinforce the rubber stock, whereas in the latter case the fillers act as extenders assisting the compounding process. Several factors have been initially recognized to determine the action of the fillers within an elastomeric matrix. These involve (i) the extend of the interface between elastomer and filler usually expressed in square meters per gram of filler, (ii) the nature of this surface evident in adsorption properties and chemical reactivity, (iii) the shape of filler particles involving aggregates and agglomerates, defining the so-called ‘structure’ and (iv) the porosity of the filler particles [4]. The latter characteristic is structurally classified into micro-, meso- and macro-porous with pore sizes of less than 2 nm, between 2 and 50 nm and larger than 50 nm, respectively [5]. The selection of the appropriate filler to reinforce a specific rubber requires a number of considerations. A high specific surface area along with the respective loading determines the effective contact area between filler and rubber [6].
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This constitutes a requisite for enhanced reinforcement emanated by the increased occluded rubber. Additionally, the type and number of the active sites on a filler surface plays a crucial role on the degree of bonding with the elastomeric matrix. In this respect treatment with coupling agents is able to modify accordingly the surface energy of the related filler [7, 8]. The induced rubber– filler interactions restrain the mobility of the macromolecular chains in the vicinity of the latter. This conformational entropy loss of the macromolecular chain is compensated by its enthalpy gain. For uncured compounds the created ‘bound-rubber’ remains un-dissolved during an extraction process in a good solvent [9, 10]. On the other hand, in vulcanizates, that immobilized rubber is characterized as rubber shell. Its presence results in an interphase creation with discernible properties [11]. Two or more fillers might come close enough so that a joint rubber shell is created. Usually, a filler network formed via joint shell construction is less rigid than that formed by direct contact of fillers. The forces induced by surface charges of fillers are able to give rise to agglomeration. Rubber trapped within a filler network would behave more as filler than as a matrix, at least in terms of mechanical properties. Under circumstances, in reinforced vulcanizates, overlaps among the three types of immobilized rubber can be in action i.e. part of occluded rubber may also be trapped rubber or rubber shell and vice versa [11]. The agglomerates found in a rubber matrix comprising of several small stable units called aggregates. These latter are made of a number of homogeneous primary entities tight together. For example, CB has a primary particle with cross-sectional dimension of 5–100 nm, whereas its aggregates consist of multiple primary particles and present dimensions of 100–500 nm. Condensing a plethora of aggregates into agglomerates a cross-sectional dimension of 1–40 lm is obtained [12]. The term ‘structure’ used in the literature for CB describes the arrangement of aggregates while the term ‘secondary structure’ is connected with the conformation of its agglomerates [3]. Nowadays, the fillers mixed with elastomeric matrices have variety of sizes and dimensions. These might include spherical-like (i.e. 0D) [13] and rod-like forms (i.e. 1D) [14–16] to platelet-like (i.e. 2D) [17] and other complex configurations (i.e. 3D) [18, 19]. Therefore, their ‘structure’ and eventually, their ‘secondary structure’ are diversified accordingly. Crucial role on the dispersion of the fillers within the elastomeric matrix plays the character of the rubber itself. Polar or non-polar elastomeric matrices must be chosen with caution so that to fit to the surface activity of the respective filler, and vice versa. The quality of rubber–filler interphase determines in great extent the properties of the vulcanizate. Its response to external stimuli is differentiated as the vulcanizate operates from low temperatures (i.e. glassy behaviour) to temperatures above the glass transition (i.e. rubber-like behaviour) [20]. At the same time, the relaxation response of the overall reinforced elastomer is altered, accordingly. The methods developed to monitor the relaxation performance of the vulcanizate supply valuable information for the determination of its structure-property relationships.
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2 Relaxation Analysis Methods 2.1 Dynamic Mechanical The type and extend of the segmental motion of the macromolecular chain in elastomeric nanocomposites along with its relaxation phenomena is linked to the physical properties of the vulcanizates. The restrained mobility of the rubber in the vicinity of filler alters its response to an applied force. This alteration is able to be closely examined when the force is periodic, induces limited deformation and does not cause any failure or appreciable fatigue [21]. Additionally, applying a sinusoidal stress and performing a temperature or frequency sweep across the desirable range, the induced strain will also alternate sinusoidal, however, presenting a lag (i.e. out of phase). This viscoelastic performance of the vulcanizate is susceptible to the various motions activated within the temperature or frequency range tested. Appropriate technique to monitor such phenomena has proved to be dynamic mechanical analysis (DMA). By this method parameters like the ability of the vulcanizate to store energy (i.e. storage modulus), to lose energy (i.e. loss modulus) and the ratio of these effects (i.e. loss factor) are plausible determined [22]. The latter dimensionless parameter is a measure of dynamic hysteresis of the vulcanizate and it is generally characterized as damping. DMA testing might be associated with different types of fixtures suitable for tensile, flexure or compression modes. The tensile mode is usually in action. It should be mentioned at this point that the dynamic properties of rubber vulcanizates investigated by a sinusoidal strain dynamic tester, which involve low to moderate deformations able to destroy the created filler network (i.e. Payne effect), are not considered in the current analysis [23, 24]. Examining the thermal transitions in elastomers, changes in free volume or relaxation times supply information for the various segmental motions. Making use of the crankshaft model, as the free volume increases, the ability of the chain segments to move in various directions also increase. Usually at very low temperatures, the localized bond movement of bending and stretching is ascribed as gamma (c) transition, which may also involve associations with water. As the temperature and the free volume increases, whole side chains and localized groups of four to eight backbone atoms are in action. Such transitions are called beta (b) transitions. The large scale motions of the amorphous regions as heating increases further are related to the glass transition [22]. Above that threshold the rubbery state appears. The modulus in that region is proportional to the number of cross-links. In elastomers, glass transition temperature often defines the lower limit of the operating temperature. For DMA testing, a frequency value close to the real world application should be chosen. Due to the fact that this appears in most of the cases to be inconvenient for the instrumentation, a set of the same parameters for all samples in the data set is considered. Generally speaking, a shift of the temperature of a transition to lower values appears by lowering the frequency. Moreover, as the testing frequency increases, the modulus usually exhibits also an increase.
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2.2 Dielectric Spectroscopy Electrical relaxation phenomena occurring in polymer and polymer composites can be studied experimentally by means of broadband dielectric spectroscopy (BDS), which is also called dynamic electrical analysis, underlying that the method is the electrical equivalent of DMA. Both methods exert a time dependent field (electrical or mechanical) and record the material’s response, with varying parameters the frequency of the applied stimulus, the sample’s temperature (in some cases the applied pressure) and the amplitude of the exerted field. In DMA the responding entities of the material are masses, while in BDS the responding ones are permanent or induced dipoles. As it can be easily understood BDS can record faster relaxation effects (characterized by sorter relaxation times) than DMA. Broadband Dielectric Spectroscopy is a powerful technique for the investigation of physical effects occurring in polymers and polymer micro- and/or nanocomposites, such as molecular mobility, polarization, conductivity, interfacial phenomena, phase changes, polymerization, crystallization, etc. [25–28]. In our days modern and convenient devices are available in reasonable prices, facilitating thus the usage of dielectric spectroscopy in basic and applied research. The employed experimental set up (device and dielectric cell) varies according to the examined frequency range. However, an analytical description of the required experimental set ups is out of the scope of the present chapter, useful experimental information can be retrieved elsewhere [25–28]. Dielectric data are recorded under isothermal or isochronal conditions scanning the frequency of the applied field in the first case and specimen’s temperature in the second one. Data can be analysed via different formalisms that is in terms of: (a) dielectric permittivity, (b) electric modulus, (c) ac conductivity, and (d) complex impedance. Electrical phenomena present in materials can be described and studied via all four formalisms. However, under specific circumstances one of them could be proved more helpful in extracting information concerning the observed physical effects. Typically, dielectric data are expressed via the real and imaginary part of dielectric permittivity and electric modulus according to Eqs. 1 and 2, respectively. Equation 3 defines tan d which is a measure of the dissipated energy upon the stored energy at each charging cycle. e ¼ e0 je00 M ¼
1 1 e0 e00 ¼ ¼ þ j ¼ M 0 þ jM 00 e02 þ e002 e e0 je00 e02 þ e002 tan d ¼
e00 M 00 ¼ 0; e0 M
ð1Þ ð2Þ
ð3Þ
where e0 , M0 and e00 , M00 are the real and the imaginary part of dielectric permittivity and electric modulus, respectively. The electric modulus presentation offers
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advantages in the interpretation of dielectric data, since the large values of (e0 ) and (e00 ) which obscure relaxation effects in the low frequency range and at high temperatures are minimized. Further, the influence of electrode polarization becomes negligible [29–32]. Relaxation processes in polymer nanocomposites arise from both the polymer matrix and the presence of the filler. Dynamics of relaxation processes can be investigated by recording the frequency of dielectric loss peaks varying the temperature of the sample. Relaxation dynamics express the influence of temperature upon relaxation time or dielectric loss peak frequency and allow investigating the type of the distribution of relaxation times (i.e. symmetrical, non symmetrical, and superposition of symmetrical and non symmetrical) [29]. From these results valuable information for the matrix/filler interactions might be obtained [33–37].
2.3 Others Among other techniques, transient mechanical testing has been used in the literature to examine relaxation phenomena in elastomeric vulcanizates. These may involve creep, set and stress relaxation, all methods investigating the result of an applied stress or strain as a function of time. Parallel to that, these methods are used as measures for ageing characteristics, low temperature resistance and resistance to chemicals [38]. A technique which currently receives increased attention for dynamic molecular interaction processes in elastomer vulcanizates is solid state nuclear magnetic resonance (NMR) spectroscopy. Utilizing methods like magic angle spinning and cross polarization, the line broadening has been reduced increasing at the same time the NMR signal sensitivity [39]. Thus, solid state NMR might be considered of comparable quality to the traditional liquid NMR. The central NMR observable is the residual dipolar couplings characterizing local chain order. In this respect, multiple quantum NMR has been utilized for the measuring of weak residual dipolar couplings in elastomeric nanocomposites [40].
3 Relaxation Phenomena in Nanocomposites with Non-Polar Elastomeric Matrices 3.1 Dynamic Mechanical Analysis Characterization Natural rubber (NR) has traditionally received most of researchers’ attention among the non-polar elastomers. Regarding the spherical-like nanofillers, which have been mixed with NR, CB constitutes a special subcategory wherein filler agglomeration is rather desirable for real-life applications. This emanates from the
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fact that an elastomer reinforced with dispersed CB agglomerates behaves, in terms of its viscoelastic response, as its filler concentration is higher than its real volume [41, 42]. This discrepancy between theory and practice has been settled considering the occluded rubber as part of the filler thus, introducing the term of ‘effective’ filler [21, 43]. Due to broad existence of huge agglomerates in high filler loadings, CB filled elastomers usually decline from their classification as nanocomposites. Nevertheless, CB has been employed in several cases as standard filler for identifying mechanisms involved in elastomeric (nano) composites. Reduction of the CB particle size in butadiene rubber (BR) resulted in a respective decrease of the damping performance as this was revealed in a tan d versus temperature graph. The compound with the lowest damping presented also the highest bound rubber value [44]. Styrene butadiene rubber (SBR) was mixed with 40 phr CB through melt blending in various percentages of sulfur curatives. A number of calendered stocks were prepared, wherein each possessed its particular percentage of curatives. By hot-pressing the individual layers together a so-called functionally graded polymeric nanocomposite (FGPNC) was produced. The increase in crosslink density of the cured layers caused a shift of the glass transition temperature from 19 to +20 °C as detected during dynamic mechanical tests. Thus, damping proved to be a sensitive indicator of crosslinking. Additionally, comparing a uniformly dispersed polymeric nanocomposite (UDPNC) with the respective FGPNC, both retaining the same equivalent of curatives, interesting results were obtained. More specifically, due to the variation of crosslink density along the thickness, FGPNC broaden the glass transition region as observed in a tan d versus temperature graph. Moreover, its tan d peak value at the Tg was lower compared to that of UDPNC, whereas glass transition temperature remained rather unaffected [45]. Considering the loss factor as a measure of material internal losses per cycle deformation during periodical excitation, its monitoring might reveal notable information for the filler performance under wear conditions. Tire applications currently involve particulate fillers like CB and silica. In this respect, master curves of tan d at the reference temperature of 20 °C and at strain amplitude of 0.5% were produced for SBR vulcanizates filled with silica and CB in 80 and 81 phr filler loading, respectively. At the low frequency region silica produced lower tan d values compared to the CB stock, indicating better rolling performance. On the other hand, as frequency increased (1–1,000 kHz) silica yielded higher tan d values compared to CB, thus higher energy dissipation. The latter frequency range is related to the excitation frequencies involved during the antilock braking system braking phase [46]. The last decade, rod-like fillers for rubber reinforcement have been the subject of vivid research. Pre-vulcanized NR latex has been mixed with water solution of multi-walled carbon nanotubes (MWCNT) and sodium dodecyl sulfate (SDS). After stirring and sonication the mixture was cast and left to dry. The DMA investigation of the vulcanizates showed an almost linear decrease of the mechanical loss factor peak value at the glass transition temperature (Tg) versus the MWCNT weight percentage. Efficient stress transfer to the filler was provoked by the
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treatment of the MWCNTs with nitric acid, which created carboxyl functionalities on the side walls of the purified nanotubes. The significantly increased storage modulus in the rubbery state of the elastomeric nanocomposites with higher filler weight percentages was accompanied by higher tan d values compared to the neat matrix [47]. Additionally, at higher MWCNT percentages (i.e. 5.4 and 8.3 wt%) the tan d peak at Tg was shifted to lower temperatures, likely due to the increased amount of SDS along with the filler aggregation observed at these loadings. Generally speaking, the enhanced mechanical performance of elastomers reinforced with carbon nanotubes (CNT) is not always accompanied by a peakshift to higher temperatures of the Tg in a loss tangent (tan d) versus temperature graph. Considering that the existence of such peak-shift hints build-up of rubber– filler interactions, its absence raises various argumentations [48]. SBR containing 23.5 wt% bound styrene was dissolved in toluene and stirred with 0.66 wt% MWCNT along with sulfur type curatives. The purified nanotubes were sonicated in ethanol and dried prior their solution mixing with SBR. The ability of the vulcanizate to absorb energy was enhanced as shown by the increase of loss tangent peak in the DMA spectrum. Additionally, the glass transition temperature was shifted slightly to higher temperatures (i.e. 2.5 °C) compared to the neat matrix. The prominent damping characteristics of the SBR/MWCNT nanocomposites were accompanied by improved storage modulus in the glassy and rubbery state [49]. The creation of functional groups on the surface of CNTs is often requisite for mixing with non-polar elastomers in order to improve the load transfer to the filler [50]. In this respect, CNTs were first pretreated with nitric and sulfuric acid followed by treatment with hydrated silica, resorcinol and hexamethylene tetramine. Ball-milling technique was exploited for the CNT prior their solution mixing with NR. It was found that the CNT preparation process created carbonylic and hydroxylic functional groups on the filler surface. Addition of 25 phr CNT in the NR matrix restricted mobility of the NR backbone lowering the damping performance of the vulcanizate along with a shift of the Tg to higher temperatures [51]. Going a step further, introducing MWCNTs in an ionic liquid ethanol solution followed by its melt mixing with an SBR/BR blend (in 1/1 ratio) significant interfacial interactions were created. The ionic liquid 1-allyl-3methyl imidazolium chloride (AMIC) was thought on the one hand to participate in rubber vulcanization via its reactive double bond and on the other hand to be tethered on the CNT surface by cation-p interactions. This strong rubber–filler coupling was responsible for the enhanced storage modulus in the rubbery region and the decreased damping at the glass transition temperature by adding 3 phr of MWCNTs. The interaction of the AMIC with the rubber and the CNT formed a thin polymeric layer around each CNT of significantly reduced mobility. As revealed in the mechanical loss factor versus temperature spectrum, apart from the main a-relaxation peak of the matrix another smaller peak at about 80 °C was detected [52]. Supportive data to the idea of a tight shell around the CNTs were given by the clear peak-shift in the Raman spectrum towards higher wavenumbers of the G-band, which is associated with the tangential C–C stretching of the graphite-like structures.
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The interactions of the macromolecular chains with the respective substrate have been found responsible for the promotion of relaxation phenomena and the shift of the glass transition temperature [53, 54]. Thus, platelet-like nanofillers is expected to vigorously define various relaxation behaviors for elastomer nanocomposites according to their affinity with the matrix. Waxy maize starch nanocrystals with a thickness of 6–8 nm, a length of 40–60 nm and a width of 15– 30 nm has been mixed with NR latex. This type of nanofiller enhanced significantly the storage modulus of the respective nanocomposites especially in the rubbery region as recorded by means of DMA. Addition of 30 wt% of starch nanocrystals resulted strongly in a reduction of the damping performance of the vulcanizates, shifting slightly the Tg to lower temperatures [55]. Lately, a lot of research has been conducted on rubber/layered silicate nanocomposites. For this 2D filler type (referred also as clay if it is of natural origin), an elastomeric nanocomposite is obtained favorably when the individual silicate sheets are surrounded by the macromolecular chains of the rubber. This scenario includes rubber confined within the silicate galleries (cf. intercalated nanocomposite) and/or individual silicate sheets dispersed in the matrix (cf. exfoliated nanocomposite) [56]. Significant role on such preferred silicate dispersion in the rubber matrix plays the organic modification of the sheets with the appropriate intercalant [57]. The organically modified clay is generally called ‘organoclay’. The relaxation phenomena involved in intercalated structures emanate from the favorable rubberorganoclay interactions and the shielded rubber by the silicate layers. Thus, less amount of elastomer is able to be deformed under dynamic conditions reducing the amount of dissipated energy [58]. In this respect, NR melt-mixed with organically modified montmorillonite (OMMT) presented reduced damping performance and a shift of the Tg to higher temperatures in a loss tangent versus temperature graph [59, 60]. In certain cases, the confined macromolecular chains within the silicate layers were charged with the appearance of a second relaxation peak (or of shoulder-like form) positioned at higher temperatures in a tan d versus temperature spectrum. Such behavior has been reported for NR/fluorohectorite nanocomposites produced via the latex route [61] and NR melt-mixed with 10 phr montmorillonite (MMT) modified with octadecylamine in tensile mode at a frequency of 10 Hz [62]. The presence of the intercalants on the surface of the clay layers should additionally affect the dynamics of the rubber chains in their vicinity. In order to amplify this behavior, SBR and ethylene-propylene-diene rubber, were filled with OMMT up to 150 phr. In both highly filled rubber/OMMT nanocomposites, a transition around 60–70 °C was detected by means of DMA. The relaxation peak raised at that temperature region in a tan d versus temperature graph was connected with the thermally induced transition of the surfactant (i.e. alkyl chains) from solid-like to liquid-like state [63]. Enhancement of the interactions of non-polar rubbers with organoclay layers through coupling agents was found to be a convenient tool for increased mechanical performance of said nanocomposites [64, 65]. SBR latex was mixed with a 2 wt% aqueous suspension of MMT and 3-aminopropyltrimethoxysilane followed by cocoagulation in 2% dilute sulfuric acid solution. The dried compound was meltblended with the sulfuric curatives and 2 phr bis(triethoxysilyl-propyl)tetrasulfide
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(i.e. Si69) prior vulcanization. The coupling of Si69 through its ethoxysilyl-propyl and sulfide groups with SBR and the Si–OH of the nanoclay decreased the mobility of the rubber chains near the platelets. Thus, the Tg of SBR/clay compounds was shifted from -49.8 to -34.8 °C when the clay layers were micro- and nano-scale dispersed, respectively [66]. In exfoliated structures of rubber/organoclay nanocomposites, the single platelets dispersed within the matrix are accompanied by their tethered intercalants. While the first restrain the rubber chain mobility through the network created [67], the latter act as plasticizer in the nanocomposite [58]. Thus, a moderate peakshift of the glass transition temperature due to the above-mentioned interplay is expected to be revealed in a loss factor versus temperature spectrum [68]. Nowadays, demanding applications of elastomers might require a combined action of fillers [69, 70]. In such cases, the relaxation phenomena involved in relevant vulcanizates are related to each filler component, as well as, to the reinforcing filler-blend efficiency. In this respect, SBR was melt-mixed with CB and organoclay. The DMA spectra of the vulcanizates reinforced with CB and OMMT both in 20 phr loading resulted in significantly increased storage modulus and decreased damping compared to the vulcanizates filled only with CB or OMMT [71].
3.2 Dielectric Probing Dielectric relaxations are related to dipolar effects, thus non-polar rubbers appear to be inactive dielectrically. Despite the technological importance of NR, which was stated in the previous paragraph, scarce studies on its dielectric response have been reported in the literature [32, 36, 59]. NR is weakly active dielectrically because of its non-polar nature. The absence of polar side groups in the polymer chain of NR is responsible for lacking of any secondary relaxation process in its dielectric spectra. Thus, the only present relaxation process in its spectra refers to the glass to rubber transition of NR (a-mode) [36]. This is located around -64 °C as it has been detected by means of differential scanning calorimetry [36, 72]. The corresponding loss peak recorded at low temperatures appears to be weak and is detected also in the NR based nanocomposites [36]. Figure 1 depicts the dielectric response of NR via the variation of the imaginary part of electric modulus (M00 ) as a function of both temperature and frequency. The dielectric response of NR has been examined in the temperature range from -80 to 50 °C and in the frequency range of the field from 0.1 Hz to 10 MHz. The presence of a single relaxation process is evident in Fig. 1. Electric modulus loss maximum (attributed to glass to rubber transition) shifts, with increasing frequency, from the vicinity of -60 to almost 0 °C. Figure 2 presents the dielectric response of a NR nanocomposite. NR has been reinforced by 10 phr (parts per hundred of NR) of layered silicates (LS). The employed LS exhibit very high aspect ratio (viz.[1,000). Intergallery distance of LS, as determined by X-ray diffraction (XRD) scattering, was found to be 0.95 nm. In the case of the nanocomposite system the basal spacing expands to 1.19 nm, indicating intercalation of
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α-mode 50 0 -50
-2 10
0 10
2 10
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equ
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y[ enc
6 10
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Fig. 1 Imaginary part of electric modulus versus frequency and temperature for pure NR
10-3 10-4 10-6
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0 10
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Fig. 2 Imaginary part of electric modulus versus frequency and temperature for NR/LS nanocomposite in 10 phr filler loading
LS by latex components [36, 72, 73]. As it can be seen in Fig. 2 the NR/LS nanocomposite spectra resemble to the ones of the pure matrix, including the primary relaxation process of a-mode. However, it should be noted that in a composite system it is expected to detect interfacial polarization (IP). IP also known as Maxwell–Wagner–Sillars (MWS) effect occurs in heterogeneous and complex
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systems, because of the accumulation of mobile charges at the interface of the constituents. At the interfaces large dipoles are formed. The relaxation of these dipoles under the influence of the applied field give rise to a slow relaxation process (characterized by enhanced relaxation time), which in most of the cases is recorded at relatively high temperatures and low frequencies [29, 74–77]. In the case of the NR/LS nanocomposite, this peak can be found in the low frequency region of its dielectric spectra in Fig. 2. At temperatures higher than 20 °C the values of (M00 ) for both systems increase rapidly likely due to the softening of the matrix. It is well known that the frequency-temperature superposition shifts dielectric loss peaks to higher frequencies with raising temperature. Relaxation mechanisms differ in their peak shift rate, reflecting the type of the involved dipoles, the influence of the molecular environment and the effect (if any) from the filler. The relaxation dynamics can be expressed in terms of the temperature dependence of dielectric loss peak frequency position, or in terms of the temperature dependence of the relaxation time. In the second case, the distribution of relaxation times can also be investigated via suitable distribution functions, which are available in the literature [26, 29]. Fast relaxation processes, such as local motions of polar side groups and small segments of the polymer chain (b- and c-mode respectively) typically follow Arrhenius type temperature dependence. IP relaxation exhibits the same type temperature dependence, although it is recorded at the opposite edge of the time scale compared to local motions. Arrhenius behaviour is expressed by Eq. 4: EA fpeak ¼ f0 exp ; ð4Þ kB T where fpeak is the frequency of the dielectric loss peak, f0 pre-exponential factor, EA the activation energy of the process, kB the Boltzmann constant and T temperature (in Kelvin). The relaxation process related to glass to rubber transition, is not characterized by a constant loss peak shift rate, and thus its temperature dependence cannot be described via Eq. 4. The temperature dependence of a-mode follows the Vogel–Fulcher–Tamann (VFT) equation, which considers that relaxation rate increases rapidly at lower temperatures because of the reduction of the free volume. VFT relation is expressed by Eq. 5: AT0 fpeak ¼ f0 exp ; ð5Þ T T0 where fpeak is the frequency of the dielectric loss peak, f0 pre-exponential factor, A a constant being the measure of activation energy, T temperature (in Kelvin) and T0 the Vogel temperature or ideal glass transition temperature. Figure 3 presents the temperature dependence of a-mode for pure NR and NR reinforced with 10 phr LS, while the values of parameters A and T0 evaluated via fitting data with Eq. 5 are listed in Table 1. The increase of parameter A in the case of the nanocomposite indicates an enhancement of the required activation energy for the onset of the process, suggesting indirectly that the presence of nanofiller exerts some restrictions on macromolecules rearrangement.
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Fig. 3 Loss peak position as a function of the reciprocal temperature of a-relaxation process, for the NR and NR reinforced with 10 phr LS systems. Lines are produced by fitting data via Eq. 5
Table 1 Fitting parameters of Eq. 5 for the glass to rubber relaxation process of NR and NR reinforced with 10 phr LS Sample A T0 (K) NR NR ? 10 phr LS
25.7 35.0
128 120
Recently, the effect of type and amount of LS loading upon the dielectric response of NR nanocomposite was studied in detail [32]. Pure clay and an organically modified type were used for the production of NR/LS nanocomposites (vulcanized and non-vulcanized versions) [32]. IP or MWS effect is present in all nanocomposites. Glass to rubber relaxation process, detected in all systems, slows down its dynamics when the NR matrix is vulcanized. Further, in vulcanized systems the temperature dependence of the dielectric loss peak position appears to be unaffected by the type of the employed LS. Relaxation spectra of non-vulcanized systems include normal-mode relaxation, in which the end-to-end polarization vector characterizes the translational motion of the whole chain. This process vanishes after vulcanization, because of the formation of cross-links [32] and becomes not detectable. An additional slow relaxation mode was observed in all NR/organoclay systems. Its origin was attributed to partially immobilized polymer chains or restricted segmental mobility at the interface layer around clay particles [32].
4 Relaxation Phenomena in Nanocomposites with Polar Elastomeric Matrices 4.1 Dynamic Mechanical Analysis Characterization In several cases elastomeric nanocomposite manufacturing requires increased rubber polarity for optimized filler dispersion in the respective matrix. However,
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tailoring the interface quality and performance in polar elastomeric nanocomposites is not an easy task. While precipitated silica dispersion in nitrile rubber is favored by increasing acrylonitrile content, a negligent effect might appear for CB reinforcement [44]. Nano-sized silica has been melt-mixed with carboxylated nitrile rubber (XNBR) at a filler loading up to 50 phr. DMA investigations in the torsion mode at a frequency of 1 Hz revealed a slight reduction of the damping performance of the nanocomposites along with a marginal peak-shift of the Tg to lower temperatures [78]. This seems to be related to the silica agglomeration and de-wetting phenomena at the respective interface. In order to overcome flocculation of silica within the rubber matrix, the first is synthesized in-situ during compounding. Silica particles were synthesized from N-(2amino-ethyl)-3-aminopropyltrimethoxysilane in XNBR. The rubber type used was containing 6.4% of carboxyl groups and 26.3% of acrylonitrile. As the amount of aminosilane varied in the rubber recipe from 5 to 20 phr, in order to produce the set of reinforced vulcanizates, crosslink density was accordingly changed. The ionic nature of the crosslinks between the carboxyl groups of the XNBR and the amino groups of the silane restrained the mobility of the macromolecular chains at the interface. This was reflected in the tan d versus temperature graph wherein apart from the arelaxation at about 0 °C, which was almost constant for all vulcanizates, another one at higher temperature (&40 °C) appeared, characterized as a0 -relaxation. As the aminosilane amount in the vulcanizates increased, the loss tangent value at the glass transition temperature decreased along with an intensity increase of the a0 relaxation [79]. The emergence of slower segmental dynamics than in the bulk rubber have been also referred for hybrid polyurethane (PUR) reinforced with polyhedral oligomeric silsesquioxanes [80]. For the same polar rubber, when it was mixed with 10 phr alumina nanoparticles of various sizes via the latex route and left to dry, decrease of the particle size of the filler resulted in a reduction of the intensity of the loss factor peak at the glass transition temperature. At the same time, a marginal tan d peak-shift to higher temperatures was detected for the compounds with small particle diameter [37]. It seems that by enhancing rubber– filler interactions less macromolecular chains contribute to a-relaxation due to their immobilization in the vicinity of the filler. As shown in Fig. 4, the alumina nanoparticles (being 25 nm in diameter) resulted in a significant enhancement of the storage modulus at the rubbery state along with a shift of the Tg to higher temperatures compared to the neat matrix. In addition, the damping of the nanocomposite was strongly reduced. Rubber–filler interface enhancement for polar elastomers has also been investigated in case of rod-like nanofillers [81]. For MWCNT compounded with hydrogenated carboxylated acrylonitrile butadiene rubber by melt blending, the presence of polar groups had small effect on the glass transition temperature [82]. It seems that significantly slower segmental dynamics require apart from the adequate rubber polarity a properly modified filler surface, and vice versa. Nevertheless, enhancement of the mechanical performance is generally reported. For platelet-like fillers mixed with polar rubbers, extended research has been conducted during the last decade. Epoxidized natural rubber (ENR) has been used
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Fig. 4 Storage modulus and mechanical loss factor of PUR and PUR/alumina nanocomposite (in 10 phr filler loading) versus temperature (in tension mode at 10 Hz frequency). Note that the line represents the neat PUR whereas line plus symbol the PUR reinforced with nanoparticles of 25 nm in diameter
as compatibilizer for NR mixed with MMT modified by octadecyltrimethylamine. Increasing the degree of epoxidization from 25 to 50 mol% (designated as ENR 25 and ENR 50, respectively) resulted in a pronounced immobilization of rubber chains in the vicinity of the platelets as this was corroborated by attenuated damping behavior [83]. Decreased damping and shift of the Tg to higher temperatures has been also reported for the same system mixed with 50 phr CB or 30 phr silica [84]. However, melt blending of only 2 phr organoclay with NR compatibilized by 10 phr of ENR 50 resulted in pronounced intercalated structures and improved strength. Moreover, in a loss factor versus temperature graph, apart from the a-relaxation at about -45 °C another peak at about -5 °C was raised. The latter was connected with the intercalated rubber populations due to the increased polarity of ENR 50 [83]. Similar segmental dynamics were marked also when the organoclay loading increased up to 10 phr [85]. The ability of ENR 50 to penetrate within silicate layers has been elucidated when it was melt mixed directly with organoclay [86]. As it is shown in Fig. 5, addition of 10 phr organoclay (MMT modified by octadecylamine) enhanced the storage modulus of the vulcanizate compared to a silica filled compound of same loading. Additionally, the a-relaxation at about -5 °C was reduced in intensity followed by a ‘shoulder’ positioned at about 10 °C, due to the segmental motions of a part of ENR 50 matrix activated at a higher temperature. Dynamic mechanical analysis do not always detect the discernible relaxation phenomena due to the intercalated rubber populations in polar rubbers reinforced with platelet-like nanofillers [87]. Nevertheless, a tan d maximum value decrease is usually observed in the relevant DMA spectra [88–90]. As the interaction of the clay intercalant with the rubber matrix increases the damping reduction might become more pronounced. In such an example, hydrogenated nitrile rubber (HNBR), which was reinforced with MMT modified with various intercalants resulted in a decreased intensity of the tan d peak at the glass transition temperature for all nanocompounds. However, when a methyltallow-bis(2-hydroxyethyl) quaternary ammonium salt was in action, the induced hydrogen bonds restrained
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Fig. 5 Storage modulus and mechanical loss factor of ENR 50 vulcanizates mixed with silica (line curve) or organoclay (line plus symbol curve) in 10 phr loading versus temperature (in tension mode at 10 Hz frequency)
Fig. 6 Storage modulus and mechanical loss factor of neat PUR/NR blend (line curve) and reinforced with 10 phr sodium fluorohectorite (line plus symbol curve) versus temperature (in tension mode at 10 Hz frequency)
further the mobility of the nitrile rubber, thus, an even lower tan d peak intensity was measured [91]. Rubber blends are usually in action in order to overcome cost and property related issues. In such an example, PUR rubber was mixed with a quite nonexpensive pre-vulcanized NR via the latex route followed by the addition of sodium fluorohectorite (FHT) slurry in 10 phr filler loading [73]. While the first rubber bears polar groups able to create enhanced interface with FHT, the latter apolar one interacts weak with the FHT layers. The characteristics of the dried compounds were resolved by means of DMA. As shown in Fig. 6, two distinct a-relaxation peaks rise at about -55 and 5 °C for NR and PUR rubber, respectively, revealing the immiscibility among these two rubbers. Addition of FHT in the compound shifted the glass transition temperature of PUR at about 10 °C, while the Tg for NR remained rather unaffected. On the other hand, the increased tan d values at the rubbery region for the nanocomposite suggest a more efficient energy dissipation mechanism than that of the plain blend. This is likely due to the ‘house of cards’ structure built by the dispersed FHT platelets, which are actually
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expelled by the pre-vulcanized NR latex particles [73]. Nevertheless, the nanodispersion of FHT platelets in the blend enhanced significantly the storage modulus of the compound in the whole temperature range.
4.2 Dielectric Probing Polar rubbers are, as expected, dielectrically active materials. Their response include contributions arising from glass to rubber transition (a-mode), local motions of polar side groups (b-mode) and from the rearrangement of small segments of the polymer main chain (c-mode). In cases where additives and/or plasticizer are present, as well as in semi-crystalline rubbers IP could be detected. The dielectric response of nanocomposites based on polar matrices compromises dipolar events arising from the matrix as well as from the presence of the filler. Further, the presence of the nanofiller might affect relaxation phenomena emanating from the polymer matrix. Synthetic PUR is a typical polar rubber used for the development of many nanocomposites. In most of the studies the employed nanofillers are alumina particles, LS and CNT. Slow and fast relaxation processes can be observed in the dielectric spectra of PUR. The latter compromise local motions of small parts of the main chain, like the ‘crankshaft’ type motion of the [CH2]x sequence (c-mode), and re-orientation of polar side groups of the main chain (b-mode), while the former are a-mode (glass to rubber transition) and IP [36, 72]. The PUR macromolecular backbone consists of hard and soft segments [36, 75, 76]. At the interface between soft and hard regions unbounded charges, arising from the stage of samples preparation, are concentrated forming large induced dipoles, which participate to the polarization effect of the system at the low frequency and high temperature regions. PUR/NR blends are also used as base for the development of rubber nanocomposites. Figures 7 and 8 show the variation of the imaginary part of electric modulus as a function of temperature and frequency for PUR and PUR/NR systems. All four aforementioned relaxations can be easily detected in the PUR spectra, while blend’s response contains also the contribution from the a-mode of NR. Reasons for choosing the electric modulus formalism have been, briefly, referred in Sect. 2.2. Relaxations of a PUR mixed with 10 phr alumina nanoparticles are depicted in the spectra of Fig. 9. Qualitatively the picture refers to the response of PUR. The presence of nanofiller within a polymer matrix, could affect the matrix glass transition temperature. For favorable polymer–filler interactions at the nanoscale Tg shifts to higher value, compared to that of the pure bulk polymer [92, 93]. On the other hand, repulsive interactions at the interface between matrix and filler, as well as the existence of pore or voids because of poor wetting level of the inclusions, or insufficient adhesion between polymer and nanofiller result in enhancement of the cooperative mobility of the macromolecules, and thus in decreasing glass transition temperature [94, 95]. Another reason, which could influence Tg in rubber/LS nanocomposites is their structural configuration [33, 36].
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10-2 10-3
Modulus''
10-1
IP
10-4
γ-mode α-mode β-mode 2 10
4 10
y[
c 0 en 0 -2 1 equ r 10 F
8 0 6 1 10
] Hz
-150
-100
50
0
-50
Temperatu
re [°C]
100
Fig. 7 Imaginary part of electric modulus versus frequency and temperature for pure PUR
10-2 10-4
10-3
Modulus''
10-1
IP
β,γ-modes of PUR
α-mode of PUR α-mode of NR
2 10
4 10
y[
c 0 en 0 -2 1 equ r 10 F
8 0 6 1 10
] Hz
-150
-100
-50
Temperatu
0
50
100
re [°C]
Fig. 8 Imaginary part of electric modulus versus frequency and temperature for PUR/NR blend
In the case of intercalated structures rubber molecules lying within the LS galleries, are able to move faster than bulk polymer chains. Confining of polymer macromolecules in a narrow space at the nanoscale level, results in isolated macromolecules, which can move or relax faster or easier, and by thus shift glass transition temperature to lower values [33]. Glass transition temperature variation
G. C. Psarras and K. G. Gatos
10-1
108
10-3 10-4
Modulus''
10-2
IP
10-5
γ-mode α-mode β-mode 2 10
4 10
y[
c 0 en 0 -2 1 equ r 10 F
8 0 6 1 10
] Hz
-150
-100
-50
Temperatu
0
50
100
re [°C]
Fig. 9 Imaginary part of electric modulus versus frequency and temperature for PUR reinforced with 10 phr alumina particles. The mean diameter of alumina particles is 25 nm
can be detected and studied by means of BDS. In the vicinity of Tg macromolecules acquire sufficient energy to increase their mobility and to relax under the influence of the applied alternating electric field. In nanocomposites shifting the loss peak of a-mode, under isothermal conditions, to higher frequency compared to that of pure bulk polymer, signifies lowering of Tg. This is quite reasonable, since shifting the loss maximum to higher frequency denotes decreasing of relaxation time, due to the facilitation of the process. On the contrary, when loss peak position of a-mode in a nanocomposite is shifted to lower frequency, with respect to the loss position of bulk polymer, Tg value enhances. It should be noted that BDS can be used not only for qualitative studying the modification of glass transition temperature in nanocomposites, but also in studying the molecular relaxation dynamics, and in identifying the type of the occurring interactions at the system’s interface. Secondary relaxations (i.e. b- and c-mode) are weak, sometimes suffer by experimental noise, and characterized by short relaxation time, and thus recorded at low temperatures and high frequencies. In some cases they mutually superimposed and/or are not easily detectable. Usually their loss peak position remains unaffected by the presence of nanofiller. The dielectric permittivity loss maximum could increase in rubber nanocomposites indicating higher consumption of energy at each charging cycle. Enhanced energy consumption could be attributed to restrictions, exerted by nanoinclusions, to local motions and reorientations [36, 72]. Figure 10 depicts the variation of real (e0 ) and imaginary part (e00 ) of dielectric permittivity with frequency, in PUR and PUR/alumina nanocomposite systems at -100 °C. Nanocomposites exhibit increased values of both parts of dielectric
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10
PUR PUR+10phr alumina (220nm) PUR+10phr alumina ( 90nm) PUR+10phr alumina ( 25nm) PUR+40phr alumina (220nm)
9
0.20
0.16
8 0.12
e ′′
e′
7 6
0.08
5 4
0.04 3 2
-1
0
1
2
3
4
5
6
0.00
logf Fig. 10 Real and imaginary part of dielectric permittivity as a function of frequency, at -100 °C, for PUR and PUR/alumina systems, varying the mean particle size and/or the alumina content
permittivity, which in the case of the system with alumina particles of 25 nm mean size attain maximum values. As responsible for this behaviour could be considered the enhancement of interfacial area with diminishing the mean alumina’s particle diameter. Additionally, the selection of the amount and the type of the employed nanoinclusions offer the possibility to tailor the dielectric response of the nanosystems. Hydrogenated nitrile rubber, commercially introduced in the 1980s, addresses advanced requirements of the automobile industry because of its excellent heat and oil resistance properties combined with superior mechanical performance [96]. HNBR is a polar elastomer wherein its acrylonitrile content and its crystallinity may vary accordingly [96, 97]. HNBR/CNT nanocomposites are reported to improve, further, the mechanical properties of HNBR vulcanizates, as well as, their wear performance [98]. The electrically conductive character of the CNTs offers an additional advantage in applications, where moving or rotating members, fabricated with HNBR, are present [99]. On the surface of moving or rotating insulating members electrostatic charges can be easily developed leading to undesirable sparks, which can cause an early failure of the component. A possible solution to the problem is to incorporate within the insulating matrix conductive inclusions, which will not lower the mechanical performance of the system and at the same time they will be able to create conductive paths, inside the composite systems, through which leakage current could flow. HNBR/CNT nanocomposite appear to satisfy these requirements, since CNTs act as reinforcing phase to both mechanical and electrical behaviour. Recently, the electrical response of HNBR and HNBR/CNT nanocomposites was examined thoroughly [99]. In HNBR and HNBR/CNT systems with low concentration of CNTs four distinct relaxation processes were detected. These are IP at the interface of crystalline and amorphous
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10-1
regions of HNBR (the existence of crystalline phase was confirmed via XRD spectra) and/or at the interface between of HNBR and CNTs, glass to rubber transition of the amorphous part of HNBR, rearrangement of polar side groups, such as –CN, and local motions of small segments of the main elastomer chain [99]. At higher CNTs concentrations, above a critical value also known as percolation threshold, the electrical response of the system varies significantly, because it switches from insulating to conductive performance. Relaxation processes vanish or notably suppressed, while the electrical behaviour of the system is dominated by the increased values of electrical conductivity [99]. The dielectric response of PUR/LS and PUR/NR/LS nanocomposites (in 10 phr filler loading) is presented in Figs. 11 and 12 respectively. Although, the observed relaxations are those mentioned previously, some modifications are present in the spectra of nanocomposites. The IP loss peak of the PUR/LS system appears to be broader and its maximum is suppressed to lower values compared to the corresponding of PUR (Fig. 7), PUR/NR (Fig. 8) and PUR/NR/LS (Fig. 12). The occurring modification has been connected to variations of the morphology of the examined systems [36, 72]. Previous morphological studies, via transmission electron microscopy and XRD [72, 73] on the PUR reinforced with 10 phr LS system, revealed the existence of two intercalated populations with different interlayer distances, namely 1.23 and 1.73 nm. Note that the basal spacing of LS powder was 0.95 nm. Further, the presence of isolated silicate layers due to partial exfoliation cannot be excluded. Under this point of view, interfaces with varying geometrical characteristics co-exist within the nanosystem, contributing to interfacial relaxation phenomena with different dynamics or relaxation times. The
Modulus''
10-3
10-2
IP
10-4
α-mode β-mode
γ-mode
2 10
4 10
y[
c 0 en 0 -2 1 equ 10 Fr
8 0 6 1 10
] Hz
-150
-100
0
-50
Temperatu re [°
50
C]
Fig. 11 Imaginary part of electric modulus versus frequency and temperature for PUR/LS nanocomposite (in 10 phr filler loading)
111
10-1
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Modulus''
10-3
10-2
IP
10-4
α-mode of PUR β, γ-modes of PUR
α-mode of NR
2 10
4 10
y[
c 0 en 0 qu -2 1 e r 10 F
8 0 6 1 10
] Hz
-150
-100
0
-50
Temperatu
50
re [°C]
Fig. 12 Imaginary part of electric modulus versus frequency and temperature for PUR/NR/LS nanocomposite (in 10 phr filler loading)
recorded broad peak can be considered as the superposition of all interfacial effects. Interfacial relaxation phenomena prominent in the low frequency range and relatively high temperatures are characterized by high values of both real and imaginary part of dielectric permittivity [31, 100, 101]. Further increase of the intensity of interfacial effects results in even higher values of (e0 ) and (e00 ). However, in the electric modulus formalism the increase of intensity of the IP process is demonstrated by reduced values of (M0 ) and (M00 ), because of its definition via Eq. 2. Thus, the recorded lower values of the modulus loss index (M00 ) constitute a strong indication for the existence of pronounced interfacial phenomena in the PUR/LS system. In all other systems, the formed IP loss peaks resemble to each other. The PUR contribution seems to dominate the spectra of all systems. Morphological inspection saw that the PUR intercalation of LS is considerably better than the one of NR [73]. This effect can be assigned to the higher polarity of PUR compared to NR, which favors the compatibility with LS [73]. Since, NR and PUR are not compatible (amplified by the prevulcanization of NR), and the intercalation of LS by NR is poor, silicate stacks are preferentially located in the PUR regions [73]. The amplitude of the IP loss peak in the case of the PUR/ NR/LS nanocomposite attains lower value compared to that of PUR/NR and PUR systems, implying the relative enhancement of IP due to the increase of the system’s heterogeneity. The temperature dependence of the dielectric loss peak position for a-mode and IP of the systems: (i) pure PUR, (ii) PUR/LS, (iii) PUR/NR, and (iv) PUR/NR/LS is shown in Fig. 13. Note that filler loading was set in 10 phr. As it can be seen IP follows an Arrhenius type behaviour, described by Eq. 4, while a-mode’s peak
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Fig. 13 Loss peak position as a function of the reciprocal temperature for the IP and a-relaxation processes, for PUR, PUR/LS, PUR/NR, and PUR/NR/LS systems (LS is incorporated in 10 phr in the respective compounds). Lines are produced by fitting data via Eqs. 4 and 5
Table 2 Activation energy for the IP process and fitting parameters of Eq. 5 for the glass to rubber relaxation process of the PUR based systems Sample IP-process a-Relaxation process
PUR PUR/NR PUR ? 10 phr LS PUR/NR ? 10 phr LS
EA (eV)
A
T0 (K)
1.29 1.37 0.60 1.29
8.3 5.8 33.0 2.4
215 225 165 246
shift is in agreement with VFT relation, as expressed by Eq. 5. Activation energy values and values of parameters A and T0 calculated via Eqs. 4 and 5 are listed in Table 2. The increase of parameter A in the case of the PUR/LS nanocomposite indicates an enhancement of the required activation energy for the onset of the process, suggesting indirectly that the presence of nanofiller exerts some restrictions on macromolecules rearrangement. Moreover the same system exhibits the lowest value of activation energy for IP denoting that its morphological characteristics, mentioned previously, facilitate the process.
5 Summary Dynamic mechanical analysis and BDS are two effective experimental techniques for studying rubber nanocomposites. Their spectra not only reveal the mechanical and dielectric behaviour of the under test systems, but also provide valuable information regarding molecular dynamics, interfacial effects, filler/matrix interactions, and phase changes increasing the impact of our knowledge upon structure–property relationships. DMA and BDS are basically complementary methods, although they are able to describe relaxations of the same type (primary and
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secondary modes). Relaxation spectra compromise contributions originating from glass to rubber transition, local motions of side groups or small side branches, rearrangement of small segments of the rubber main chain, IP at heterogeneous systems and phase changes. DMA is a well established method for the determination of glass transition temperature, secondary relaxations of side groups, and the investigation of mechanical energy storage capacitance and damping characteristics of rubber nanocomposites. BDS besides glass to rubber transition is able to follow faster molecular relaxation processes arising from reorientation of polar side groups and small chain segments, IP effects due to the accumulation of unbounded charges at the interface of the constituents and conductivity phenomena. Further, qualitative information for the type of the occurring interactions between nanofiller and elastomeric matrix can be revealed.
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Modeling and Simulation of Polymeric Nanocomposite Processing Teik-Cheng Lim
Abstract As a consequence of its large surface area to volume ratio, the embedment of nano-particles in polymers leads to large interface area and hence considerable interphase volume region between the nanoparticles and the polymer. The resulting bulk properties in the context of nano-particle filled polymers therefore differ from the use of conventional particle reinforcement. While the mechanical, electrical and other solid state physical properties of polymer nanocomposites can be easily obtained due to the static nature of the molecular and continuum modeling, the same cannot be said so for the case of polymer nanocomposite processing due to the dynamic flow nature inherent in the latter. This chapter lays down the common rules adopted in modeling of polymer nanocomposite processing. Beginning from the various interatomic and intermolecular potential energy functions that are indispensable for molecular modeling, the chapter presents two major approaches for molecular modeling of polymer flow. Recent results arising from the use of molecular modeling is then summarized with emphasis on the glass transition temperature and the rheological properties of the polymer melt with the presence of nano-scale fillers. The chapter concludes with the advantages of molecular modeling techniques for understanding the nanoparticle-filled polymer in the context of flow processing. Keywords Modeling
Nanocomposites Nanoparticles Processing
T.-C. Lim (&) School of Science and Technology, SIM University, 535A Clementi Road, 599490 Singapore, Republic of Singapore e-mail:
[email protected] V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_4, Ó Springer-Verlag Berlin Heidelberg 2011
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1 Introduction As the phrase suggests, the processing of polymer nanocomposites involves flow of melt polymers and polymeric solutions with nano-scale fillers, i.e. fillers whereby at least one of the dimensions is equal to or less than 100 nm. As in conventional polymer composite processing—such as compression molding, extrusion, injection molding, pultrusion, resin casting, rotational molding, spinning, electrospinning, thermoforming and transfer molding— processing of polymer nanocomposites are governed by the same processing parameters and boundary conditions associated with the processing machinery and flow pathway. Unlike the conventional polymer composite processing, the unique properties arising from the large surface area to volume ratio of the nano-scale fillers, hence the large interfacial interaction between the surface area of the nano-scale fillers with the adjacent polymeric chain. The nano-scale fillers can possess aspect ratio greater than 1 such as carbon nanotubes or other nanorods [1–5], aspect ratio of 1 such as buckminsterfullerene, carbon black or other nanoparticles [6–10], and aspect ratio less than 1 such as montmorillonite clay or other nano-platelets [11–15]. Polymers that were processed with nanoscale fillers include polypropylene [16], polyimide [12, 17], elastomers [18, 19], etc. Modeling of polymer nanocomposite processing differs significantly from computer simulation of bulk polymeric melt flow with and without conventional fillers, which uses rheological properties of the polymer, the processing condition, the filler volume fraction, the filler shape (aspect ratio) and the mold geometry. The accuracy of any simulation of the structural properties of a nanomaterial, and its processing thereof, are ultimately governed by the appropriate use of interatomic and intermolecular potential energy functions used in the simulation, and the length of time-step increment imposed on the simulation. The interatomic potentials covered include (1) van der Waals systems, (2) covalent interactions and (3) ionic interactions. Computer modeling of polymeric nanocomposites can be categorized into two broad methods, namely the molecular dynamics (MD) approach and the Monte Carlo (MC) approach. The MD approach has been adopted by Douglas et al. [20, 21], Smith et al. [22–26], Keblinski et al. [27–29] and others. The MC approach was employed by Mark et al. [30–32], Mattice et al. [33–35], Vacatello [36–38] and Termonia [39], among others. This chapter begins with a review of the various interatomic potential energy functions that are used in both the MD and MC approaches. The fundamentals of MD and MC methodologies are then discussed with reference to polymer processing. Finally, results from the MD and MC techniques from various research groups are discussed with emphasis on the effect of nano-scale fillers on the rheological and other flow properties of the melt polymers with nano-scale fillers.
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2 Potential Energy Functions In molecular modeling, one aims to obtain the flow prediction and the corresponding polymer chain conformation either on the basis of minimum energy or on the basis of steepest descent of energy. The Hamiltonian H¼
N X
Vi þ
i¼1
þ
N3 X
N 1 X
N2 X U2b ri;iþ1 þ U3b hi;iþ1;iþ2
i¼1
i¼1
N1 X N X U4b si;iþ1;iþ2;iþ3 þ UNb rij
ð1Þ
i¼1 j iþ3
i¼1
gives the overall energy of the molecular system, whereby V is the kinetic energy of each atom, U is the potential energy, N being the number of atoms in the system, while the bond length, bond angle and torsional angle are denoted as r, h, and s respectively. The kinetic energy Vi ¼
pi ; 2m
ð2Þ
where pi and m refer to the momentum of ith atom and its mass, applies during polymer flow. The interatomic potentials, U, can be broadly classified into bonded and non-bonded interactions. The bonded interactions can be further categorized into 2-body (stretching) potential, 3-body (bending) potential and 4-body (torsional) potential. The non-bonded interaction takes place between two atoms not directly bonded within the same molecule as well as between atoms of different molecules, as a result of van der Waals interaction (electrodynamics) and also due to Coulombic interaction (electrostatics). The 2-body potential considers the energy stored when two adjacent atoms, bonded to one another, undergoes relative displacement. The simplest potential is based on a simple harmonic oscillator consisting of a mass m attached to a rigid wall by a spring of stiffness kr. Hence the load F applied in stretching the spring by an amount r can be written as ð3Þ
F ¼ kr r according to Hooke’s Law, and F ¼ ma ¼ m
d2r dt2
ð4Þ
according to Newton’s Second Law. Since the summation of energy, H, is constant 1 2 mv þ U ¼ H; 2
ð5Þ
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then differentiating with respect to time, with due consideration of (4), leads to Hinchliffe [40] dr dU ¼ 0: ð6Þ Fþ dt dr Since (dr/dt)=0, we have Z U ¼ Fdr:
ð7Þ
Substitution of (3) into (7) yields 2 1 U2b ¼ kr rij rij;e 2
ð8Þ
where rij is the bond length between ith and jth atom with subscript e referring to equilibrium bond length, and rij - rij,e = r [41–49]. An obvious disadvantage of this potential, also known as the harmonic potential, is that the potential energy curve forms a quadratic curve, which is symmetrical about the equilibrium bond length, thereby predicting an erroneously large energy at long distance. This drawback was overcame by the introduction of the Morse potential [50], 2 ð9Þ U2b ¼ D 1 exp a rij rij;e where D and a are the Morse parameters. These, as well as the equilibrium bond length, can be obtained by performing curve-fitting of data from ab initio results or more traditionally from spectroscopic data. This potential function has been widely used when large bond stretching is anticipated [51–55]. In most cases, the harmonic potential would be sufficient except where long polymeric chains are simulated due the significant bond stretching during molecular chain entanglement. The 3-body potential considers the molecular energy stored when the bond angle, h, formed by three atoms is changed as a consequence of relative displacement. The harmonic version of bond bending can be written in a way that resembles (8), i.e. 2 1 U3b ¼ kh hijk hijk;e 2
ð10Þ
where kh refers to the bending stiffness by the angle formed by atom i, j and k [45– 49]. A more frequently adopted alternative to (10) is the Quadratic Potential in Cosine [41–44, 51–56] 2 1 kh cos hijk cos hijk;e : 2
ð11Þ
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Equation (11) gives rise to lower energy in comparison to (10) when h [ h0 but higher when h \ h0. Furthermore, the asymmetric curve from (11) is more realistic. The 4-body potential quantifies the amount of energy stored when a series of 4 atoms undergoes torsion about its own axis. A direct way to describe the change in energy resulting from the change in torsional angle of s is the periodic function [45, 46, 48] 1 U4b ¼ ks ½1 þ s cosðnsÞ ð12Þ 2 in which ks is the torsional stiffness, n is the order of symmetry (or periodicity of the potential) and s is the barrier to rotation (or phase factor). A more frequently adopted torsional potential is written in the form [41–44] 5 X an cosn s ð13Þ U4b ¼ ks n¼0
whereby the coefficients an are obtained empirically [57] as a0 = 1, a1 = 1.31, a2 = -1.414, a3 = -0.3297, a4 = 2.828 and a5 = -3.3943. In other forms to (13), the torsional stiffness and coefficients are re-grouped such that U4b ¼
3 X
Cn cosn s
ð14Þ
n¼0
where Cn (in kJ/mol) are empirically obtained coefficients [47, 49, 52–54, 56, 58–60]. The non-bonded potential energy functions are the van der Waals interaction imposed upon atoms that are not bonded together in the same molecule, between atoms of different molecules, and also between molecules or clusters or atoms. The force between such atoms is highly repulsive when brought closer together but mildly attractive, with a minimum and followed by decay, when separated. There are two major sets of potential functions used: exponential-6 function and the Lennard-Jones (12-6) potential. The exponential-6 function is written [47, 54, 61] in the original version as ð15Þ UX6 ¼ Arij6 þ B exp Crij or in the loose form as [46] ( " 6 #) h
6 rij i f r exp f 1 UX6 ¼ 4e r f6 f 6 rij
ð16aÞ
or ( UX6 ¼ e
6 #) " 6 rij f rij;e exp f 1 rij;e rij f6 f6
ð16bÞ
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whereby A, B and C are coefficients which depends on the system under consideration and f is a scaling factor which reduces to the conventional Lennard-Jones form at long range when f = 12 and coincides with the minimum well-depth of the conventional Lennard-Jones potential when f = 13.772, with e being the dissociation energy while r is the internuclear range at which U = 0. Both e and r are also known as the LJ parameters, which are comparable to the more frequently applied Lennard-Jones function [47–49, 52–55] " 6 # r 12 r : ð17aÞ ULJ ¼ 4e rij rij or ULJ
" 6 # rij;e 12 rij;e : ¼e 2 rij rij
ð17bÞ
Since the potential energy decays at long distance, the Lennard-Jones potential is truncated (e.g. [41–44]) such that ( h 12 6 i 4e r=rij r=rij ; rij 2:5r ð18Þ ULJ ¼ 0; rij [ 2:5r because the potential energy magnitude would be less than e/60 when rij = 2.5r. Other forms of the Lennard-Jones potential energy have been attempted by Noid and Pfeffer [51] " 6 # 9 r r 3 ULJðNPÞ ¼ e 2 ð19Þ rij rij and for the case of interaction between molecules, 12 6 r r ULJðMolÞ ¼ 24e 2 13 7 : r r
ð20Þ
The Coulombic interactions (also known as electrostatic or ionic interactions) is written as Uc ¼
X q2 Aq2 ¼ 4pe0 rij 4pe0 R i6¼j
ð21Þ
whereby e0 = permittivity of free space, R = nearest-neighbor separation, rij = interionic distance = aijR, and A is the Madelung constant defined as A¼
X 1 i6¼j
aij
:
ð22Þ
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Due to the dynamical nature of the polymeric chain as well as the enormous number of atoms involved, simplifying assumptions have been introduced [41–45, 48, 51–54, 56]. Firstly the groups of CH2 and CH3 are lumped together to form united atom models, and secondly the cross terms are neglected. In some instances, bond stretching is neglected (e.g. [56]) while others neglected bond torsional energy [51]. The individual atoms in CH2 and CH3 groups has been taken into account by Noid et al. [47] and Schoen et al. [62] to check the extent of improved accuracy. It was found that modeling of CH4 and CF4 each as united atom models provide better static results whereas consideration of the 4-site bonds for these molecules results in better dynamic properties when matched against experimental results [62, 63]. As such, the united atom model is applicable to structural and materials property simulation, while consideration of all atoms is highly recommended for process modeling.
3 The Molecular Dynamics Methodology Arising from Newton’s second law, (4) and the force-potential relation, (7), we have mi
d 2 ri dU ¼ 2 dri dt
ð23Þ
The finite difference approach has been adopted to solve the equations of motion by numerical method so that the positions and velocities of each atom at time t leads to those properties at the next time step, t ? dt. The differential equation shown in (23), can be easily solved using the finite difference method. One of the most commonly used approximation is the Verlet [64] algorithm, which requires information pertaining to the position ri(t) of each atom and the corresponding atom’s acceleration ai(t), as well as the position of the atom from the previous time step, ri(t - dt). To pave a way for finite difference solution, Taylor’s expansions for ri(t ± dt) about ri(t) is performed to give 1 ri ðt þ dtÞ ¼ ri ðtÞ þ vi ðtÞdt þ ai ðtÞðdtÞ2 þ 2
ð24Þ
1 ri ðt dtÞ ¼ ri ðtÞ vi ðtÞdt þ ai ðtÞðdtÞ2 2
ð25Þ
and
respectively. Substitution of (25) into (24) leads to the Verlet algorithm: ri ðt þ dtÞ ¼ 2ri ðtÞ ri ðt dtÞ þ ai ðtÞðdtÞ2 þ
ð26Þ
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of which the accurate is valid up to the fourth order (dt)4. Subtraction of (25) from (24) gives the atomic velocity vi ðtÞ ¼
ri ðt þ dtÞ ri ðt dtÞ 2dt
ð27Þ
which is accurate to the order of (dt)2. With these equations in place, the Verlet algorithm is performed as follows. Starting with the atomic positions from two time steps, i.e. r(t) and r(t - dt), the atom’s acceleration a(t) is calculated from (4). The position of each atom at the next time step r(t ? dt) can then be calculated from (26). Due to the dynamical nature in flow processes, there is a need to obtain the kinematic energy. Hence the atomic velocity v(t) is calculated from (27). As information of r(t - dt) and r(t) leads to r(t ? dt), the information of r(t) and r(t ? dt) leads to r(t ? 2dt) in the next round of calculation. The advantages of the Verlet algorithm are that the advancement of positions is all performed in one step and conserves energy well even with relatively long time step. In addition, this technique is time-reversible, highly compact and easy to program. However, the velocities at t can be calculated only after r(t ? dt) are known. Another disadvantage is that a trajectory can only be started if information of two time steps are available, i.e. r(t) and r(t - dt) instead of just one time step, i.e. r(t) and v(t). As an alternative to (27), the Verlet algorithm was extended to the Velocity Verlet algorithm [65], which can be generally summarized as 1 vðt þ dtÞ ¼ vðtÞ þ ½aðtÞ þ aðt þ dtÞ: 2
ð28Þ
The algorithm begins with r(t) and v(t) to calculate a(t). The next position r(t ? dt) is calculated using (24). Then the mid-step velocity is calculated as dt 1 ¼ vðtÞ þ aðtÞdt: ð29Þ v tþ 2 2 followed by acceleration a(t ? dt) at the next time step. The velocity move is then completed using dt 1 vðt þ dtÞ ¼ v t þ þ aðt þ dtÞdt: ð30Þ 2 2 The main advantage of the Velocity Verlet algorithm over the original Verlet algorithm is that the former can start with positions and velocities at time t, and that the kinetic energy at time (t ? dt) is readily available. As with the original Verlet algorithm, the Velocity Verlet algorithm is also numerically stable, simple to program, time-reversible, and conserves energy well even with relatively long time steps. The only drawback of the Velocity Verlet algorithm is that it requires the calculation of two velocity steps.
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The MD approach has been applied for simulating various physical behavior of polymeric systems by numerous researchers including Rigby et al. [41–44], Noid et al. [47, 51–55] and Jin et al. [48].
4 The Monte Carlo Methodology The Monte Carlo (MC) method uses random numbers to find particle displacements during molecular simulation. A frequent way of implementing an MC calculation is to use the Metropolis method. Beginning from a starting configuration of s a random distribution of particles (atoms) and molecules, one atom i is picked randomly and displaced in a random direction by a random amount, from n rm i to ri , subject to the maximum displacement being the adjustable parameter drmax. The alteration of potential energy of the system, dUmn, as a consequence of this prescribed atomic movement is then obtained for an assumed form for the interatomic potential. Suppose the atomic displacement gives rise to decreased energy (dUmn B 0), then the new position is unconditionally accepted. However if the move leads to increase in energy (dUmn C 0), then this prescribed atomic movement is accepted only conditionally, subject to Boltzmann probability factor, exp(-dUmn/kBT). The MC method was the first simulation technique to obtain the equilibrium behavior of liquids and solids [66]. With sufficiently large number of configurations sampled, the average value of the energy fluctuations give the constant volume-specific heat in liquid and, in the case of solid, fluctuations in the stress tensor determine to the elastic constants [67]. The MC method has also been applied to bulk polymers (e.g. [68–72]). The MC technique by Vao-soongnern et al. [73] simulates the molecular structure of polyethylene as a nanofiber. As in previous works (e.g. [68]), these molecular chains are placed on a diamond lattice whereby every second site has been removed, thereby giving rise to the second nearest neighbor diamond (2nnd) lattice. Computation of the polymeric nanofiber chain consisted of two parts, firstly the generation of the initial polymer nanofiber structure, which was then followed by relaxation of the initial structure to thermodynamic equilibrium. The initial step employs only self-avoiding walks with the excluded volume condition. In the following step, both the intermolecular as well as the intramolecular potential functions are then imposed [68, 69]. The initial structure is then relaxed so as to minimize its potential energy according to the Dynamic Monte Carlo technique [69–71]. The equilibrated free-standing nano-film structure was calculated from previous work by Mattice et al. [72]. The fiber is then formed via collapsing this thin film after extending, about 3 to 4 times, on one side of the periodic box in the direction that is perpendicular to the normal axis of the thin film surface. This new box would be sufficiently large in order that no parent chain interacts with its images. As such, the periodic boundary appears to apply only in one direction for this step. Two methods for fiber formation were then introduced, namely the
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elongation method and the cut-and-press method. In the elongation method, the nanofiber is formed by generating in a similar way to that of thin film formation from bulk. In the cut-and-press method, a new nanofiber is obtained from an existing larger fiber structure. Specific details of these two methods can be found elsewhere [73]. It has been shown that when the number of chains and chain length are the same, both approaches lead to equivalent nanofiber structure. The parameters simulated include the shape of the cross-section (which is almost circular), the radial density profile, the bead distribution, the segregation of the chain ends, the local orientation, the global equilibrium properties of the chains, as well as the surface energy of the nanofiber.
5 Simulation Results A number of simulation results have been reported in the literature. In this section, the findings are reviewed, and commented with particular reference to the effect of nano-scale filler presence during polymeric processing. An MD study on the influence of control parameters on nanoparticle clustering in polymer was performed by Douglas et al. [20]. Simulated results give quantitative description in which (1) the polymer-matrix strength increases, at a decreasing rate, with the nanoparticle loading; and that (2) the polymer-matrix strength decreases with increasing temperature [20]. Investigations on the glass transition temperature of pure polymers and those with nano-scale fillers were performed by Lamm and Yann [74], Papakonstantopoulos et al. [75] and others. Lamm and Yann [74] obtained the glass transition temperature of pure polyimide (PI) as well as PI with 20.23% weight of octaaminophenyl silsesquioxanone (OAPS) and PI with 10.45% weight of octahydro silsesquioxanone (OHS), in which the transition temperature is marked by a change in slope of specific temperature plotted against temperature. The MD result reveals an increase and decrease of the glass transition temperature with the addition of OAPS and OHS respectively. The modeling approach by Papakonstantopoulos et al. [75], on the other hand, obtained the glass transition temperature by the change of slope of the polymer density against temperature. Again, a shift on the glass transition temperature is detected via MD with the inclusion of nano-fillers in the pure polymer. In modeling the viscosity g of liquid polymers, one considers the shear stress s arising from the shear strain rate (dc/dt) as S dc s¼g ð31Þ dt whereby a Newtonian flow is denoted by S = 1 while the range 0 \ S \ 1 corresponds to a Power-Law fluid. For the case of Newtonian fluid, the inclusion on an additional term that describes the yield shear stress sy denotes a Bingham plastic, i.e.
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dc s ¼ sy þ g : dt
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ð32Þ
On the other hand, the Casson fluid appears to be a blend of the Power-law fluid and Bingham plastic, and is written as rffiffiffiffiffi pffiffiffi pffiffiffiffi dc s ¼ sy þ g ð33Þ dt or rffiffiffiffiffiffiffiffiffi dc dc s ¼ sy þ 2g sy þ g2 : dt dt
ð34Þ
As can be seen in all the above fluids, the viscosity g of the fluid affects the shear stress of the flow, which in turn influences flow speed and required applied pressure for processing. The viscosity g and shear stress relaxation modulus G(t) of a liquid polymer nanocomposite can be calculated by the Einstein relations [76, 77] 2 t 32 + * XX Z V 0 0 4 Pab ðt Þdt 5 g ¼ lim ð35Þ t!1 12kB Tt a b6¼a 0
and GðtÞ ¼
V Pab ðtÞPab ð0Þ kB T
ð36Þ
where V = system volume, kB = Boltzmann’s constant, T = temperature, Pab(0) = initial value of the off-diagonal element of the stress tensor, Pab(t) instantaneous value of the off-diagonal element of the stress tensor at time t, and h i averaged value over the entire trajectory and all 6 off-diagonal elements of the stress tensor. The normalized polymer nanocomposite viscosity, as a function of nano-filler volume fraction /f is given by [26] 5 Fð/f Þ ¼ 1 þ /f þ 4:94/2f þ C/3f 2
ð37Þ
where C = constant. Plots of normalized viscosity of polymer matrix nanocomposite against the specific interfacial area between polymer and nanoparticles show that the normalized viscosity increases rapidly and weakly with the specific interfacial area when the nanoparticle and polymer is attractive and neutral respectively [26]. The same simulation also shows that when the nanoparticle and polymer is repulsive, the normalized viscosity of polymer matrix nanocomposite decreases with the specific interfacial area between polymer and nanoparticles. Using a chain length of N = 20, Douglas et al. [21] showed that (1) the nanofiller reinforced polymer, just like the pure polymer, experiences viscosity decay
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with increasing shear rate, and that (2) the viscosity of nano-filler reinforced polymer has higher viscosity than that of pure polymer by one order. While these two results are not surprising, it is striking to note that, with other parameters fixed, the nano-rod fillers give the polymer the highest viscosity, followed by icosahedra (or spherical-like) nano-fillers. The nano-platelet fillers give the lowest viscosity to the polymers [21]. Although the difference in the polymer viscosity for the use of these three nano-fillers are small compared to that of pure polymer, the simulated result nevertheless provides evidence on the influence of nano-filler shape on the overall performance of nanocomposite processing.
6 Conclusions Arising from its length scale, nano-scale structures are best modeled at the molecular level, which justifies the need to incorporate knowledge on the interaction between atoms, cluster of atoms and molecular chains. This aspect sets nano-scale modeling apart from simulation of bulk properties of conventional polymer composites, which can be predicted using composite mechanics principles. Modeling and simulation of polymer nanocomposites allow a way for visualization and analysis of the processing performance in regard to the speed of production as well as the cost or energy consumption in the processing method for various combinations of polymer melt/solution, filler geometries, processing temperature & pressure, and other conditions before a prototype processing is tested. Visualization of molecular flow enables the detection of any hazardous collision of nanoparticles on the inner surface of the tool geometry, so that the tool designer can design or select tool geometries that can ensure smoother flow so as to reduce or eliminate the generation of debris, which can influence the flow pattern of the polymer nanocomposite processing as well as the properties of the final product. The visualization also enables the detection of multi-phase flow, and the extent of this flow if any, so as to redesign the tool geometry and/or recalibrate the processing condition in order to reduce energy loss arising from the second phase flow. Developmental cost is hence effectively reduced.
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Deformation-Induced Structure Changes in Elastomeric Nanocomposites Shigeyuki Toki and Benjamin S. Hsiao
Abstract During tensile deformation, nanofillers can orient and align with polymer chains to reinforce the strength and modulus of elastomeric nanocomposites. The presence of nanofillers can also enhance the orientation of surrounding polymer chains and accelerate the strain-induced crystallization behavior. Conventionally, natural rubber, styrene-butadiene rubber and thermoplastic elastomer (ethylene-propylene copolymer and poly-urethane) are routinely reinforced with fillers. In this chapter, the behavior of nanofiller-enhanced straininduced crystallization of elastomeric nanocomposites, based on natural/synthetic rubbers and fillers of varying sizes (from microns to nanometers) such as carbon black, multi-walled carbon nanotube (MWCNT), carbon nanofiber (CNF), nanoclay (NC) during deformation, is described.
1 Introduction Elastomer is defined as the kind of soft material, which can be stretched to a large deformation and exhibits a spring-like characteristic, capable of recovering almost to the original length when released. This behavior is caused by the mobility of the long flexible polymer chains (with a low Tg) and the corresponding network structure that can receive and transfer the stress due to the chain movement (i.e., the origin of entropy modulus). The most widely studied elastomer is natural rubber (NR), which also has a very broad range of practical applications. In fact, NR is the second most consumed biopolymer on earth (cellulose is the most
S. Toki and B. S. Hsiao (&) Department of Chemistry, Stony Brook University, Stony Brook, NY 11794-3400, USA e-mail:
[email protected] V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_5, Ó Springer-Verlag Berlin Heidelberg 2011
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consumed one). NR contains naturally polymerized poly-isoprene chains as the major component, as well as proteins, carbohydrates and phospholipids. These non-rubber components also reinforce the rubber component (polyisoprene chains) as nanofillers. Carbon black is the most widely used filler to reinforce the mechanical properties of NR. During deformation, the carbon black particulates often behave as additional (physical) network points and accelerate strain-induced crystallization of the NR matrix. Recently, with the advance of nanomaterial synthesis, many newly available nanofillers have also been used to reinforce the mechanical properties of both natural and synthetic elastomers. For example, multi-walled carbon nanotubes (MWCNT) have been routinely added to styrene-butadiene rubber (SBR), NR and thermoplastic elastomers to enhance their mechanical properties. As the size of MWCNT is much smaller, the amount of MWCNT needed to achieve the same reinforcing properties is only about one-third of the amount of carbon black typically used. As MWCNT has a thread-like shape (its diameter is usually tens of nanometers but its length is in microns), its homogeneous dispersion capability in the viscous/rubbery matrix is usually a concern. For this purpose, proper surface modification of MWCNT is necessary to induce good dispersion capability. Because of the unique shape of MWCNT, its filled elastomer also showed peculiar behavior after deformation due to the complex relaxation and interactions of polymer chains and MWCNT. Carbon nanofiber (CNF) is an analog of MWCNT but with a larger diameter. It has also been used to reinforce the properties of various elastomers. The surface modification of CNF is relatively easier than that of MWCNT, thus CFN may be more convenient to use to induce strain-induced crystallization of elastomers (such as ethylene-propylene copolymer). Another class of nanofillers that has also been demonstrated to enhance the mechanical and physical properties (such as gas permeability) of elastomers is nanoclay. For example, surfactant-modified montmorillonite can also accelerate strain-induced crystallization in NR and thus increase the total crystallinity. Similar to nanofibers, nanoclays can also orient and align upon stretching, where the presence of nanoclays can also induce orientation of surrounding polymer chains and overall strain-induced crystallization, resulting in reinforcement of the mechanical properties of this class of nanocomposite.
2 Non-Rubber Components in Natural Rubber Natural rubber (NR) is composed of 94 wt% of rubber components and 6 wt% of non-rubber components such as proteins, carbohydrates, phospholipids and metal ions. Rubber components mainly consist of polyisoprene (with 100% of cis-1,4 structure) having functional groups at chain ends. The synthetic analogue of natural rubber is polyisoprene (e.g. IR2200), which is 100% polyisoprene with 98.5% of cis-1,4 structure. IR2200 showed significantly inferior mechanical properties to
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NR. Although the inferior properties have been partially attributed to the lower content of cis-1,4 structure (98.5%), the main reason is due to the absence of nonrubber components in the compound, which have been confirmed recently [1, 2]. It is now clear that in NR, non-rubber components create naturally occurring networks and multi-scaled structures simultaneously. Recently, small-angle X-ray scattering (SAXS) and small-angle neutron scattering (SANS) techniques have been used to reveal the existence of multi-scaled structures (from nanometers to micrometers) in NR [1, 2]. Figure 1 illustrates the integrated SAXS profiles of NR, IR2200 and de-proteinized natural rubber (DPNR). The comparison of these profiles indicates that the excess scattered intensity in NR is due to the scattering from proteins and other non-rubber components. The corresponding wide-angle X-ray diffraction (WAXD) patterns of NR, DPNR, Gel and Sol fractions of the NR sample (Gel and Sol are classified as the heavy part and light part of toluene dissolved NR) are shown in Fig. 2. It was seen that a small amount of large crystals is present in NR, as well as in Gel samples, but not in DPNR and Sol samples. Therefore, protein agglomerates and crystals (identified as quebrachitol), as well as phospholipid micelles, are distributed throughout the NR matrix, which can be confirmed by optical microscopy shown in Fig. 3. The decrease in the sizes of these structures by mechanical processes can be seen from the SAXS profiles of NR and NR0 samples shown in Fig. 4, where NR was the original sample and NR0 was subjected to a two-roll milling for 5 min at room temperature. It is seen that the SAXS profile of NR0 (dotted line) at the scattering vector q = 0.003 shows much lower scattering intensity than that of NR. This confirms that the large scattering structures (in micron size) are decreased by mechanical milling. On the other hand, SAXS profiles of NR0 at the scattering vector q = 0.04 show a higher scattered intensity than those of NR. This indicates that the amount of smaller scattering structures (in nanometer size) is
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Fig. 2 2D-WAXD patterns of Gel, NR, DPNR, and Sol samples under static conditions (i.e., no deformation) (Reproduced from Reference [2] with permission from John Wiley & sons Inc.)
Fig. 3 Optical micrograph (A) and cross polarized micrograph (B) of NR. Three kinds of microstructures in NR were identified: (a) protein aggregates, (b) quebrachitol crystals and (c) micelles of phospholipids and fatty acids (Reproduced from Reference [2] with permission from John Wiley & sons Inc.)
increased by mechanical milling (the above observations are indicated by two arrows in Fig. 4). In NR, the functional groups at both ends of polyisoprene chains can react with non-rubber components to form a naturally occurring network. This naturally occurring network is responsible for the high green strength and high viscous behavior of unvulcanized NR, where the effect can be clearly identified in Fig. 5 (the stress–strain curves of unvulcanized NR, DPNR and IR). It is seen that unvulcanized IR shows lower modulus and no spring-back behavior. In fact unvulcanized IR behaves as a polymer melt because it has a low Tg (-70°C) and no network structure. Unvulcanized NR shows strain-induced crystallization, which was clearly supported by the in situ WAXD measurement during extension and retraction (as illustrated in Fig. 6).
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Fig. 4 SANS (symbols) and SAXS profiles (thin lines) of nature rubber (a) before (NR) and (b) after milling (NR0) without cross-linking (Reproduced from Reference [1] with permission from American Chemical Society.)
Fig. 5 Stress–stress (elongation ratio) curves of NR, DPNR, and IR samples at room temperature. Reproduced from Reference [1] with permission from American Chemical Society.)
Based on the above studies, the non-rubber components in NR, such as protein agglomerates, quebrachitol crystals and phospholipid micelles, can be considered as nanofillers that offer additional physical crosslinking points to enhance the properties.
3 Carbon Black Filled Natural Rubber Carbon black has been used extensively to reinforce general rubber compounds such as NR and styrene-butadiene rubber (SBR). The amount of carbon black in general rubber compounds is often from 30 to 50 phr (i.e., parts per hundred parts of rubber). Carbon black is composed of primary particles (the diameter ranges from 20 to 30 nm) and aggregates (primary particles make aggregates, the radius
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of gyration of the aggregates is around 1 micron and the aspect ratio is around 5–6). These numbers vary with the kind of carbon blacks used. The presence of carbon black generally increases modulus and tensile strength however, it decreases elongation at the break of rubber compounds. The most effective means to enhance the tensile strength is to increase elongation at break through strain-induced crystallization [3]. The stress–strain relation and selected 3D WAXD patterns of carbon black-filled NR during extension and retraction are shown in Fig. 7. Each WAXD image was taken at the average strain value indicated by the arrow. It is seen that the presence of carbon black in fact accelerates the strain-induced crystallization process [4]. It is well known that the flexibility of polymer chains in NR depends mainly on the content of the cis-1,4 structure (the trans-1,4 structure is too rigid). During vulcanization, the transformation of cis- to trans- would occur. The higher curing temperature and the longer curing time can cause more transformation. Therefore, the curing condition is one of the important factors to keep the polymer chains flexible. The sulfur vulcanization process would create mono-sulfur bridges, di-sulfur bridges and multi-sulfur bridges in vulcanized compounds. To assure the chain flexibility between the network points, the choice of accelerator is also important. The choice of carbon black and the amount of carbon black also turns
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out to be an effective tool to keep the polymer chain flexible. For example, the crystal fraction of B2 (i.e., NR prepared at lower cure temperature, shorter cure time and less carbon black amounts) is higher than that of B1 (i.e., NR prepared at higher cure temperature, longer cure time and more carbon black amounts) in Fig. 8. The stress–strain curves showed that the elongation at break and tensile strength of B2 are larger than those of B1 (as illustrated in Fig. 9). It is interesting to note that the tensile strength of the B2 compound officially reached 42.5 MPa (which is the highest value ever reported [3] in the literature we searched.) under the condition of ISO-37. Therefore, the presence of carbon black accelerates straininduced crystallization and simultaneously increases the elongation at break leading to very high tensile strength.
4 Multi-Walled Carbon Nanotube Filled Elastomers Multi-walled carbon nanotube (MWCNT) has become one of the most widely investigated nanofillers in elastomeric composites. [5,6] The diameter of MWCNT
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ranges from one nanometer to tens of nanometers and the length ranges from several micrometers to millimeters or even centimeters (as illustrated in Fig. 10a–d). Figure 10a shows that MWCNT has a broad distribution of length (0.1–5 lm) and diameter (10–50 nm). Figure 10b shows the curling structure of an isolated tube. Figure 10c shows the entanglements of the tubes and the formation of an interconnected network structure. Figure 10d shows the nanostructure of a multi-walled carbon nanotube with several layers of carbon graphite and a hollow core. The dispersion of MWCNT in rubber is usually poor, as shown in Fig. 10e–h (TEM images of SBR with 4 phr of MWNTs). In these figures, a bundle (Fig. 10e) and black spots (Figs. 10f, g) are observed. The magnified view of one bundle (Fig. 10h) shows that MWCNTs are orientated. AFM images of SBR filled with 10 phr of MWCNTs (Fig. 11) showed a larger scale aggregation than that
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Fig. 10 TEM images of MWNT and corresponding composites. a–d: pure MWNT; e–h: SBR/4 phr MWNT composite (Reproduced from Reference [6] with permission from Elsevier Science)
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observed in Fig. 10, as large bright domains were seen. In AFM imaging, these bright domains can be ascribed to filler agglomerates. Stress-strain curves for SBR and MWCNT/SBR composite samples are shown in Fig. 12. The reinforcement effect by MWCNT was quite significant on mechanical properties of composites, since the presence of 1 phr of MWCNTs has increased the modulus by 45% and tensile strength by 70%. This behavior can be explained by theoretical predictions [5], which indicate that the modulus can be increased by both volume fraction and aspect ratio of the fillers. Figure 13 shows the comparison of experimental results and theoretical predictions using the Guth model and Halpin/Tsai model, respectively. To fit the experimental data, the aspect ratios of the fillers for both models were chosen to be 40 and 45,
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respectively. The results suggest that MWCNTs in SBR were aggregated since the real aspect ratio of each MWCNT particle is more than 100 or even 1000. The behaviors of filler aggregates during deformation were also observed by AFM, where typical results are shown in Fig. 14. In this figure, tensile deformation of the sample resulted in orientation of the filler aggregates, which appeared as the
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white regions. At the highest stretch ratio, the bundle aggregate was found to break up into long straight structures. After stretching and retraction, the SBR compound filled with 10 phr MWNT exhibited a large degree of permanent orientation, where the stress–strain relation is shown in Fig. 15. The hysteresis or stress-softening (Mullins effect) behavior could be explained by the loss of elastic chains taking place at the polymer-filler interface. It was found that the reinforcement effect of MWCNT in NR is almost the same as discussed above, which has also been reported in the literature [6]. MWCNT filled polyurethane elastomer showed peculiar behavior because of the complicated interaction between MWCNT and polymer crystals [7]. The concept of the peculiar response of MWCNT filled polyurethane was that both MWCNTs and strain-induced crystals of soft segments in polyurethane can act as physical cross-linkers, controlling the mechanical and stimuli responsive behavior of this class of composites.
5 Carbon Nanofiber Filled Elastomer Carbon nanofiber (CNF) is a tubular structure with the sidewalls composed of angled graphite sheets, whose configuration is quite different from carbon nanotube (CNT). Chemically modified CNFs (MCNFs) have been used to reinforce ethylene-propylene rubber (EPR) [8]. In the modification scheme, the surface of MCNFs was grafted with short polypropylene chains, resulting in good dispersion in the EPR matrix. The chosen CNFs had a diameter ranging from 60 to 150 nm and a length ranging from 30 to 100 lm. The EPR has 84.3 mol% of propylene content co-polymerized with ethylene monomer using a metallocene catalyst. The polypropylene segments would crystallize and act as network points, whereby MCNF can increase the crystal content of EPR as nuclear agents for crystallization of the PP segments.
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Fig. 16 Stress–strain curves and selected WAXD patterns acquired during stretching of pure copolymer (unfilled circles) and 10 wt% nanocomposite (filled circles) at room temperature (Reproduced from Reference [9] with permission from Elsevier Science)
The stress–strain curves and selected WAXD patterns acquired during stretching of the unfilled EPR and 10 wt% MCNF/EPR nanocomposite samples at room temperature are compared in Fig. 16. It was found that the WAXD patterns exhibited differences between filled and unfilled samples, where the presence of an outer ring (at a scattering vector q = 18.4 nm-1) was seen in the filled sample but not in the unfilled sample. This scattering peak corresponds to the 3.4 Å spacing, which is inter-shell spacing within the CNF structure (d002). The circularly averaged WAXD intensity profiles for both filled (10 wt% MCFN) and unfilled EPR samples taken at different strains are shown in Fig. 17. It was seen that during deformation, crystal structure transformation from c to a phase occurred. (We note that the crystal morphology of isotactic polypropylene usually contains a phase lamellae with epitaxial growth of c phase). A small amount of c phase crystals was found to remain in the filled samples, but the corresponding amount of the c phase in the unfilled samples was lower. Apparently, the presence of MCNF affected the process of destruction of c phase crystal during deformation. The stress–strain relation of unfilled and 10 wt% MCFN filled EPR elastomers taken at 55°C are shown in Fig. 18. It is clear that the tensile strength, elongation at break and toughness of the 10 wt% MCNF filled samples were significantly higher than the unfilled sample. Selected integrated WAXD profiles of filled and un-filled samples taken at different strains are shown in Fig. 19.
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Fig. 17 Deconvolution of circularly averaged WAXD intensity profiles into amorphous and crystalline components. The WAXD patterns were acquired during stretching at room temperature of pure copolymer (a–c) and 10 wt% nanocomposite (d–f). The corresponding strains are indicated (Reproduced from Reference [9] with permission from Elsevier Science)
It was found that the higher initial crystallinity and higher orientation of polymer chains in the filled sample could lead to thicker crystals that melt at higher temperatures. The load experienced by the polymer matrix is effectively reduced by the network formation of MCNF fillers. The reduced load effect became more efficient at higher temperatures. Therefore, the strong bonding in the MCNFpolymer interface is essential for the mechanical reinforcement of nanofiller filled elastomeric composites. The efficient stress transfer in the MCNF/EPR composite can be attributed to two factors (1) MCFN is an effective crystal nucleating agent for crystallization of the matrix and (2) the presence of MCNF
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Fig. 18 Stress–strain curves and selected WAXD patterns acquired during stretching of the pure copolymer (unfilled circles) and 10 wt% nanocomposite (filled circles) at 55°C (Reproduced from Reference [9] with permission from Elsevier Science)
forms an additional network structure that enhances the stress transfer during deformation.
6 Nano-Clay Filled Natural Rubber Nano-clay filled NR usually does not contain exfoliated clay layers, rather it contains finely dispersed clay tactoids (i.e., nano-clay stacks containing around 10 layers) in the NR matrix. These tactoids are largely isolated and separated from each other without forming a filler framework. The average thickness of the clay tactoids is 100–500 nm, but the average lateral size is considerably larger. The presence of clay tactoids can result in a faster crystallization rate and different morphology of NR upon deformation. This is because the large interfacial surface areas introduced by nano-clays facilitate the overall chain orientation during deformation, resulting in an increase in total crystalline fraction. The experimental observations are described as follows. Figure 20 shows TEM images and SAXS patterns of a chosen nano-clay/NR nanocomposite. The TEM image indicated the tactoids were finely dispersed and separated at 10–40 nm apart. Directional SAXS patterns (A: face-on view; B and C: edge-on views) showed that nano-clay stacks were aligned parallel to the film plane with preferred orientation due to processing. Figure 21 shows stress–strain relations during stretching and retraction for unfilled and three filled (NR/Na+-MMT, NR/O-MMT, and NR/O-LAP)
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Fig. 19 Deconvolution of circularly averaged WAXD intensity profiles into amorphous and crystalline components. The patterns were acquired during stretching at 55°C of the pure copolymer (a, b) and the 10 wt% nanocomposites (c, d). The strain was zero for (a) and (c), 2.8 for (b) and 22.7 for (d) (Reproduced from Reference [9] with permission from Elsevier Science)
Fig. 20 Representative TEM images (I: scale bar 1000 nm; II: scale bar 50 nm) and SAXS patterns (A: face-on view; B, C: edge-on views) of the NR-NC1 nanocomposite (Reproduced from Reference [10] with permission from American Chemical Society)
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Fig. 21 Stress–strain curves and selected synchrotron WAXD patterns during extension and retraction cycles for the chosen samples. The corresponding stress–strain curve for the NR/ O-LAP nanocomposites is included for comparison. 2D WAXD patterns for (A) the unfilled NR, (B) NR/Na+-MMT, and (C) NR/O-MMT are illustrated at selected strains (Reproduced from Reference [10] with permission from American Chemical Society)
nanocomposites where Na+-MMT represents sodium montmorillonite, O-MMT is organically modified montmorillonite and O-LAP is modified synthetic hectrite laponite clay). Selected WAXD patterns of these samples showed that nano-clays increased strain-induced crystallization and decreased the onset strain of straininduced crystallization. Figure 22 shows the evolution of crystallinity index (CI) with strain, which exhibited significant hysteresis behavior as in the stress–strain relations in Fig. 21. Contrary to the stress values during the extension-retraction cycle, the CI values are larger during retraction than those during extension. The onset strain of crystallization was determined by the interception of the regression line in the plot of CI against the strain in Fig. 22. The onset strains of all these nano-clay composites are smaller than those of unfilled NR. The difference in the aspect ratio of these nano-clays might play an important role in the onset strain and the maximum CI of these composites. At the stretching ratio a \ 3, the orientation and alignment of highly anisotropic nano-clay particles occurred; while at 3 \ a \ 4,
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Fig. 22 Crystallinity index (CI) as a function of strain during the stretch-recovery cycle for (a) unfilled NR, (b) NR/Na+-MMT, (c) NR/OMMT, and (d) NR/O-LAP samples. The solid line is only a guide to the eye (Reproduced from Reference [11] with permission from American Chemical Society) Un-stretched state
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Fig. 23 Schematic model of deformation state of nano-clay filled natural rubber (Reproduced from Reference [11] with permission from American Chemical Society)
these highly anisotropic nano-clay particles are completely aligned along the deformation direction forming a physical network. This physical network would favor the alignment of the NR chains resulting in the increase of crystallization rate.
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The presence of interfacial adhesion between nano-particles and rubber matrix could induce an early promotion and enhancement of overall crystallization of NR chains under uniaxial stretching. Based on the above results, a schematic model of the deformation process in nano-clay filled NR is shown in Fig. 23.
7 Conclusions Based on recent studies, the inclusion of nano-fillers in the elastomeric matrix can effectively generate another level of physical crosslinking network. Under deformation, the presence of nanofiller network can accelerate the orientation of surrounding polymer chains and increase the strain-induced crystallization behavior; it can also enhance the stress transfer efficiency of the composite materials. Without question, the role of nanofillers to enhance the mechanical and physical properties of elastomers will become increasingly important as more versatile and functional nanofillers are generated in the future. Acknowledgements The financial support of this work was provided by the National Science Foundation (DMR-0405432) and Bridgestone.
References 1. Karino, T., Ikeda, Y., Yasuda, Y., Kohjiya, S., Shibayama, M.: Non-uniformity in natural rubber as revealed by small-angle neutron scattering, small-angle X-ray scattering, and atomic force microscopy. Biomacromolecules 8, 693 (2007) 2. Toki, S., Hsiao, S.H., Amnuaypornsri, S., Sakdapipanich, J., Tanaka, Y.: Multi-scaled microstructures in natural rubber characterized by synchrotron X-ray scattering and optical microscopy. J. Polym. Sci. Polym. Phys. 46, 2456 (2008) 3. Toki, S., Hsiao, B.S., Amnuaypornsri, S., Sakdapipanich, J.: New insights into the relationship between network structure and strain-induced crystallization in un-vulcanized and vulcanized natural rubber by synchrotron X-ray diffraction. Polymer 50, 2142–2148 (2009) 4. Toki, S., Minouchi, N., Sics, I., Hsiao, B.S., Kohjiya, S.: Synchrotron X-ray scattering: tensile strength and strain-induced crystallization in carbon black filled natural rubber. Kautschuk Gummi Kunststoffe 61, 85 (2008) 5. Tosaka, M., Kawakami, D., Senoo, K., Kohjiya, S., Ikeda, Y., Toki, S., Hsiao, B.S.: Crystallization and stress relaxation in highly stretched samples of natural rubber and its synthetic analogue. Macromolecules 39, 5100–5105 (2006) 6. Bokobza, L.: Multiwall carbon nanotube elastomeric composites: a review. Polymer 48, 4907–4920 (2007) 7. Atieh, M.A., Girun, N., Mahdi, E., Tahir, H., Guan, C.T., Alkhatib, M.F., Ahmadun, F.R., Baik, D.R.: Fullerenes. Nanotubes Carbon Nanostruct. 14, 641–649 (2006) 8. Koerner, H., Price, G., Pearce, N.A., Alexander, M., Vaia, R.A.: Remotely actuated polymer nanocomposites—stress-recovery of carbon-nanotube-filled thermoplastic elastomers. Nat. Mater. 3, 115–120 (2004)
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9. Kelarakis, A., Yoon, K., Sics, I., Somani, R.H., Hsiao, B.S., Chu, B.: Uniaxial deformation of an elastomer nanocomposite containing modified carbon nanofibers by in situ synchrotron Xray diffraction. Polymer 46, 5103–5117 (2005) 10. Carretero-Gonzalez, J., Verdejo, R., Toki, S., Hsiao, B.S., Giannelis, E.P., López-Manchado, M.A.: Real time crystallization of organoclay nanoparticle filled natural rubber under stretching. Macromolecules 41, 2295–2298 (2008) 11. Carretero-Gonzalez, J., Retsos, H., Verdejo, R., Toki, S., Hsiao, B.S., Giannelis, E.P., Lopez-Manchado, M.A.: Effect of nanoclay on natural rubber microstructure. Macromolecules 41, 6763–6772 (2008)
Thermally Stable and Flame Retardant Elastomeric Nanocomposites O. Cerin, G. Fontaine, S. Duquesne and S. Bourbigot
Abstract This chapter is dedicated to thermally stable and flame retardant elastomeric composites. Two approaches are considered: the synthesis of elastomeric nanocomposites, where the nanoparticles are dispersed at the nanoscale, and the incorporation of nanofillers at high loadings where agglomerate of nanoparticles are observed in the elastomeric matrix. The chapter is mainly focused on the key parameter influencing the flame retardancy, that is to say the dispersion state, and the ways to improve it by the modification of the matrix and/or of the nanofiller. A particular attention is also paid to the creation of synergies between different types of nanoparticles or between nanoparticles and conventional flame retardant additives.
1 Introduction Elastomeric materials are used in a plethora of application fields: packaging, aerospace and transports, electrical devices, etc. Some of these fields are directly concerned by fire regulations. Indeed, the inherent flammability of elastomers can O. Cerin, G. Fontaine, S. Duquesne and S. Bourbigot Univ Lille Nord de France, 59000 Lille, France O. Cerin, G. Fontaine, S. Duquesne and S. Bourbigot ENSCL, ISP-UMET, 59652 Villeneuve d’Ascq, France O. Cerin, G. Fontaine, S. Duquesne and S. Bourbigot USTL, ISP-UMET, 59655 Villeneuve d’Ascq, France O. Cerin, G. Fontaine, S. Duquesne and S. Bourbigot (&) CNRS, UMR 8207, 59652 Villeneuve d’Ascq, France e-mail:
[email protected] V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_6, Springer-Verlag Berlin Heidelberg 2011
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be the source of fire hazards. Important safety aspects in case of fire should permit an early detection of fire and should find efficient means to fight it, but the best way to guarantee the safety of people is to provide a non-flammable or at least a material exhibiting low flammability. Various methods can be used to protect elastomers or in general polymeric materials from fire. It can be cited the graft-polymerization consisting in inserting into the polymeric backbone flame-retardant monomers [1] or more simply the blending of the matrix with a polymer exhibiting higher thermal resistance properties [2]. The most commonly used approach to make thermally stable or low-flammable materials is the incorporation in the polymer of flame retardant particles, micro- or nano-dispersed in the matrix [3]. This method is generally preferred to the others because it is easily compatible with industrial processes and because it offers a good compromise between economical considerations, mechanical, thermal and fire properties. Moreover, there is a growing interest upon nanoparticles because they can be used in small quantities (a few percent) providing the same performance and thus limiting the degradation of the mechanical properties of the resulting material. This chapter is organized in five parts. First the basics of flame retardancy will be commented in order to provide the reader the flame retardancy strategies, and in particular the mode of action of nanoparticles. In the second part the fire behaviour of the different elastomer classes and the ways to improve it will be investigated. The distinction between the nanocomposites, were a real nanodispersion is achieved, and the elastomers containing nanoparticles will be highlighted. The flame retardancy induced by nanoparticles in elastomeric nanocomposites will then be presented, and in another part the role of the incorporation of nanofillers in elastomeric matrices in terms of reaction to fire will also be commented. The potential future advances for designing flame retardant elastomers will be discussed in the last part.
2 Basics of Flame Retardancy To understand the various strategies to make flame retardant materials, a synthetic overview of the basics of flame retardancy will be presented in the following. At first, why is a polymer flammable? When it is exposed to heat or flame, during a fire scenario for example, the temperature of a polymer increases, leading to its thermal degradation. Thus chemical bonds of the polymeric chains are broken to generate highly flammable volatiles. These volatile compounds spontaneously form combustible mixtures with air which ignite easily and burn with a high velocity. So, a low thermal stability associated to the release of highly flammable volatile molecules is responsible for the flammability of the material. Second, how is it possible to decrease flammability? Two major ways can be followed: influencing physically the combustion process, or leading to a chemical
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action in the condensed and/or in the gas phase. Among the processes dedicated to the condensed phase can be cited: • the artificial cooling of the matrix, by the addition of fillers which decomposes endothermically during combustion (metal hydroxides, carbonates, etc.), thus delaying the polymer degradation [4]; • the creation of a protective layer. The additives can form during combustion a shield with a low thermal conductivity which can reduce heat transfer. It then reduces the degradation rate of the polymer and decreases the release rate of the pyrolysis gases. This is the basic principle of the intumescence process. Intumescence is defined by the creation during combustion of an expanded foamed cellular charred layer on the surface which protects the underlying material slowing down heat and mass transfer [5]. The gas phase modifications concern: • the dilution of the flame-feeding gases by non-flammable ones. For example the use of metal hydroxides releasing water, or ammonium compounds generating ammonia [6]; • the flame inhibition. The radical mechanism of the combustion process occurring in the gas phase is interrupted by the radicals generated by the flame retardants. Halogenated and some phosphorus-based compounds are involved in such reactions. The various ways presented above give an overview of the flame retardant principles. Nevertheless the described processes generally not occur singly but are part of a complex process in which many modifications occur simultaneously with one dominating. Third, what are the common flame retardant nanoparticles and what is their influence on the flame retardancy of nanocomposites? Some categories of nano-additives are widely used to enhance the polymers properties, among them the thermal stability and/or the flame retardancy. The most common ones are layered silicates (clays) and silica, metal hydrates and oxides (Fe2O3, TiO2), calcium carbonates and graphene-based nanofillers. More exotically can be used polyhedral silsesquioxane (POSS) or layered double hydroxides LDH. The mechanisms of action of the nanoparticles in elastomeric composites are commented in the following. Clay minerals are part of the larger class of silicate minerals. Included in this layered silicates family are the natural montmorillonites (MMT). Undoubtedly montmorillonite clay is the most commonly used silicate for producing thermal resistant and flame retardant nanocomposites. According to Gilman et al. [7], the flame retardant mechanism of clay based nanocomposites is a build-up of a high performance carbonaceous silicate char on the surface during burning. Two hypotheses can explain how MMT-rich char surface is produced [8]. There is an assumption that the MMT is precipitated during pyrolysis as a result of progressive gasification of the polymer, but another approach suggested a migration of MMT during and after decomposition of the nanocomposite structure due to the lower
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surface energy of MMT [9]. This layer insulates the underlying material and slows the mass loss rate of decomposition products. Recently Leszczynska et al. [10] studied the mechanism that controls thermal stability of polymer/MMT nanocomposites. According to these authors, there are several factors that influence thermal stability, among which: • a labyrinth effect induced by intercalated or exfoliated structure of MMT, which limits oxygen diffusion inside the nanocomposite sample, • a steric effect, due to the interactions between the polar MMT layers and the polymer matrix thus limiting the polymer chain motion, • a barrier effect, which protects the bulk of sample from heat, decreases the rate of mass loss during thermal degradation of polymer nanocomposite and generates more intensive char formation on the surface, • a catalytic effect of the clay effectively promoting char-forming reactions. Apart from the layered silicates, another class of nanofillers is able to impart thermal stability or flame resistance through the creation of a physical protective barrier: the graphene-based nanoparticles such as graphite, carbon nanotubes and carbon nanofibres. Carbon nanotubes can be seen as elongated fullerenes: they are graphene sheets rolled up into a hollow cylinder, with each end capped with half of a fullerene molecule. There are two major types of carbon nanotubes: singlewalled carbon nanotubes (SWNTs) and multi-walled carbon nanotubes (MWNTs). Carbon nanofibres CNFs are also composed of graphene particles, stacked together and forming fibre-like structures. These particles act principally in forming a charred residue, playing a protective role during combustion, as shown in poly(methylmethacrylate) by Kashiwagi et al. [11]. Concerning graphite nanoplatelets, two types can be mentioned: natural graphite (NG), and expanded graphite (EG). The effective method of preparing the expanded graphite is by rapidly heating the pre-treated natural graphite to a high temperature, which separates the platelets and also to some extent destroys the crystalline structure. This treatment generally allows a better dispersion of the platelets. Their mode of action is quite similar to that of carbon nanotubes [12]. The most popular fillers among the nano-metal hydrates are nano-aluminium trihydrate (ATH) Al(OH)3 and nano-magnesium dihydroxide (MDH) Mg(OH)2 because of their high level of flame retardant performance associated to a low cost. Nevertheless a relative high amount of filler (typically 60 wt% in thermoplastics) is needed to achieve sufficient flame retardant properties [13]. Magnesium dihydroxide Mg(OH)2 acts in decomposing endothermically at relative low temperature (353C) into magnesium oxide and water, thus releasing non combustible gas (water steam). This process dilutes the concentration of any other gaseous products and also decreases the concentration of fuel available for combustion. The decomposition also generates the oxide residue MgO, creating a protective barrier that has relatively high heat capacity, reducing the amount of thermal energy available to further degrade the substrate. Acting in a similar way as MDH, ATH releases energy from the flames by decomposing into aluminium oxide and water, and creates an alumina barrier. Another particle exhibits comparable behaviour:
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the nano-layered double hydroxides (LDH). Its structure is based on brucite (Mg(OH)2)-like layers. The flame retardation mechanism of LDH has been studied in some matrices. According to Jiao et al. [14] they act in EVA as a combination of aluminium trioxide and magnesium dihydroxide. LDH decomposes and absorbs the heat in different temperature range. However this interpretation of the mode of action of LDH is probably not complete: according to Zhang et al. [15] nano-LDHs have a capacity to catalytically oxidize a char-rich compound in the presence of O2; they can accelerate the carbonization of the char-rich compound to obtain a more sustainable char, improve the graphitization degree of the formed char, and further encourage the formation of an intumescent and heat-insulating carbonaceous layer to some extent. This is confirmed by the work of Zammarano et al. [16, 17] who highlighted that some modifiers (i.e. sulfonate anions) catalyze charring reactions in LDH based nanocomposites during thermal degradation, enhancing the formation of a carbonaceous char and decreasing the release rate of combustible volatiles. We have seen that nanocomposites exhibit similar mechanisms of flame retardancy, that is to say the creation of a protective layer hindering the combustion of polymer. Three parameters are considered crucial to achieve good flame retardancy: the nanodispersion state, the resistance and the formation rate of the protective barrier and the viscosity of the melt [18]. Particular attention has to be paid to the morphology of nanocomposites. Indeed, the filler must be nanodispersed to talk about nanocomposites, and this dispersion is a key parameter of flame retardancy. Finally, which parameters describe fire behaviour or thermal stability? The thermal stability of materials is generally characterized by thermogravimetric analyses (TGA). The mass loss of the material is recorded as a function of a temperature ramp. Two types of thermal degradation can be implemented: nonoxidative decomposition, under nitrogen or another inert gas, and the oxidative decomposition under air. Fire behaviour can be described through three major parameters: the ignitability, the contribution to flame spread and the heat release. Moreover side effects can be observed, the most important one is the emission of smoke during combustion. Depending on the material application fields some specific tests can be implemented to simulate the desired conditions, for example the Federal Aviation Administration (FAA) seat cushion flammability test, which characterizes the involvement of aircraft seat cushions in cabin fires. More generally, three tests are usually used since each of them quantifies the ignitability, the contribution to flame spread and/or the heat release. These tests are the Limiting Oxygen Index (LOI), the UL-94 test and the cone calorimeter test. LOI is determined according to ISO 4589 [19] and ASTM D2863 [20]. It consists in determining the concentration of oxygen which will just support combustion. Another small-scaled test described in UL-94 [21] determines the tendency of a sample either to extinguish or to spread the flame once the specimen has been ignited. There are twelve classifications specified in UL-94 that are assigned to materials based on their behaviour
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regarding burning, flame propagation and dripping. The most common ones are V-0 to V-2 classifications (V-0 considered the best) for the vertical burning test and HB for the horizontal one. Cone calorimeter tests [22, 23] allow the simulation of the conditions of a fire in a small bench scale test. The measured parameters are among others the Rate of Heat Release (RHR), the Total Heat Release (THR) and the Time To Ignition (TTI). These parameters permit to evaluate the contribution to a fire of the material (reaction to fire).
3 Fire Behaviour of Elastomers Depending on their own thermal stability and reaction to fire, the different elastomers do not need the same level of protection against fire. There exist various types of elastomers. Can be cited the hydrocarbon rubbers, and among them, some thermoplastic and polyolefin elastomers. The thermoplastic rubbers include polyisoprene/natural rubber (NR), thermoplastic polyurethane (TPU) and derived flexible foams, ethylene-vinylacetate copolymer (EVA), polybutadiene (BR), block copolymers such as styrene-butadiene-styrene (SBS). The thermoplastic polyolefins assign poly(octene-ethylene) elastomer (POE) and the ethylene-propylene-diene monomer (EPDM). All these polymers present a high flammability (LOI of 17 vol% for NR [24] and 18 vol% for SBS as an example) as they thermally decompose into smaller hydrocarbon molecules, these molecules forming high fuel-value gases. As an example, PU foams are known to be great contributors to fire hazard [25]. Other types of elastomers, the inorganic or semi-inorganic rubbers, have the advantage to present intrinsically high thermal resistance and low flammability. Indeed, replacing some hydrocarbon part by an inorganic one results in a polymer with higher heat resistance. Compared to hydrocarbons, the high inorganic content of these elastomers have lower fuel value, thus reducing the flammability. These polymers are fluoroelastomers, silicones and chloroprene. Silicones, in particular polydimethylsiloxanes (PDMS) exhibit an intrinsically flame retardant behaviour with a peak of RHR (PRHR) comprised between 60 and 150 kW/m2 [26]. Lyon et al. [27] reported the development of some brand-new materials: polysilphenylene-siloxane and polyphosphazene elastomers, specially designed to present intrinsically high fire resistance (Fig. 1).
Fig. 1 Chemical structure of polysilphenylene-siloxane and polyphosphazene
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Polysilphenylene-siloxanes are ultra thermally stable elastomers. Compared to PDMS their thermal decomposition occurs 200C later, that is to say beyond 500C. Concerning polyphosphazenes, their high nitrogen and phosphorus content is responsible for their flame retardant properties. Indeed, the cone calorimeter tests done at an external heat flux of 50 kW/m2 revealed the heat release rate of a polyphosphazene is four times lower that current aircraft polyurethane elastomer. The better flame retardancy is attributed to a gas phase mechanism, where phosphorus compounds in a low oxidation state inhibits the radicals present in the flame. However, higher smoke production (five times more) is induced by the presence of phosphorus. The flammability of an elastomer depends thus on its chemical composition. So, a way to modify its combustion behaviour would be to chemically modify it. Different strategies are reported here: the design of new thermally stable and flame retardant elastomers (such as polyphosphazene) and the modifications of the polymer backbone by grafting. In their review, Levchik et al. [28] report the modifications of the structure of thermoplastic polyurethane (TPU) leading to better flame retardant properties. Can be cited the incorporation of a phosphorus-based molecules in the pendant groups [29] or in skeleton [30]. Polyurethanes with the phosphorus in the pendant groups were prepared by N-alkylation of a linear PU. LOI data showed that the fire resistance of TPU is slightly enhanced. To incorporate phosphorus in the skeleton some phosphorus-based monomers were copolymerized with 2-hydroxyethylmethacrylate to form a hydroxy-containing copolymer, used as a polyol in the synthesis of PU. The so-modified PUs were characterized through TGA, which revealed their earlier degradation temperature but concurrently their higher char yield. Najafi-Mohajeri et al. [31] synthesized a series of ferrocene-modified polyurethane block copolymers and evaluated their thermal stability and reaction to fire through TGA, LOI and cone calorimetry. It was shown that a small amount of ferrocene reduced the PHRR by 40–80% depending on the ferrocene type and content, but the LOI was not significantly modified. The interest of the incorporation in the backbone of the flame retardant is also demonstrated, since only 0.51% ferrocene in the backbone structure reduce the PRHR by 60% whereas PU containing ferrocene at 15% as an additive showed a reduction of only 54% compared to unmodified PU. Levchik et al. [28] also report the copolymerization of TPU and silicones. An amine-functionalized siloxane incorporated at 5 wt% in the TPU chain allows a PRHR reduction of 79% compared to the virgin polymer. In comparison a nonfunctionalized siloxane used as an additive at 5 wt% leads to a PRHR decrease of 70%. The chemical modification of the polymer backbone allows the synthesis of thermally stable and/or flame resistant elastomers. Up to now, the examples presented in this chapter concerned raw polymers, with no mention to nanocomposites. However, it is noteworthy that the ways described in this part could be perfectly applicable to elastomeric nanocomposites, since the decreased
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flammability of the matrix imparts better flame resistance to the whole material. Moreover, nanoparticles themselves can provide some specific fire properties to elastomer nanocomposites.
4 Flame Retardant Nanocomposites Very few scientific literature deals with elastomeric nanocomposites and flame retardancy. Indeed, up to now, only 17 articles and 9 patents were published on this subject (Database SciFinder, keywords ‘‘elastomeric nanocomposites’’ and ‘‘flame retardancy’’, March 2010). Moreover, the aspect of nanodispersion when investigating flame retardancy is not always commented and is often assumed when nanoparticles are incorporated in polymeric matrices without further characterization. Since we know it is a primordial aspect, we will make the distinction between the nanocomposites, were a real nanodispersion is achieved, and the elastomers containing nanoparticles. The following part focuses on nanodispersed flame retardant particles in elastomers, and on the parameters controlling this dispersion.
4.1 Nanofillers in Nanocomposites Some nanofillers can induce the formation of a residue during degradation, this residue acting as flame retardant by a barrier effect. The two major additive families playing this role in elastomers are the layered silicates and the graphenebased additives.
4.1.1 Layered Silicates The efficiency of clays is linked to their structure modifications. By introducing trivalent aluminium and iron ions into the clay silica sheets, or bivalent magnesium and iron ions into the central sheet, the sheet package is negatively charged, charge which is compensated by the presence of cations such as Na+ (for raw clay) between the individual sheet packages. These cations can be easily replaced by others having a stronger binding affinity with polymer matrices. Organic cations are used in order to improve the compatibility between the organic polymer matrix and the montmorillonite layers. Montmorillonite containing these organic cations are called organo-modified montmorillonite. Some of these organo-modified clays were dispersed at 2.5 wt% at different degrees in elastomeric PU by in-situ polymerization [32]. The organic cation used as a modifier in the clay structure is methylbis-2-hydroxyethyltallow alkyl quaternary ammonium, presented in Fig. 2.
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Fig. 2 Quaternary ammonium salt used as a modifier in montmorillonite
The investigation of the fire behaviour of the nanocomposite was carried out through UL-94 test (3.2 mm thickness). The clay nanocomposite allows the suppression of the dripping phenomenon during the test. Char develops quickly upon ignition; both the char layer and reduced mobility of polymer chains prevent the flow and dripping caused by polymer melting. Nevertheless, the nanocomposite remains non-classified. Montmorillonite clays (organo-modified or not) through their action on the rheology of the nanocomposites towards temperature and their char-forming catalytic effect, allow the increase of the thermal stability of the materials and sometimes of their flame retardant properties. Some other nanoparticles, among the layered silicates, can be proposed as potential flame retardants as they are supposed to provide similar properties. Can be cited talc, mica (known to increase thermal stability of PDMS [33]) or kaolin.
4.1.2 Graphene-based Nanofillers The incorporation of multi-walled carbon nanotubes was evaluated in ethyleneoctene copolymer POE [34]. Thermogravimetric analyses revealed that thermal decomposition of the polymer matrix was retarded in the MWNTs nanocomposites. This result may be attributed to a physical barrier effect, resulting from the fact that MWNTs would prevent the transport of decomposition products of the polymer nanocomposites. The use of such nanoparticles was also reported in elastomeric EVA (40% vinyl acetate) [35]. Multi-walled carbon nanotubes MWNT and carbon nanofibres CNF at 4 wt% were incorporated by solution blending. The thermogravimetric analyses of the obtained materials revealed a heat stability increase with MWNT (decreased degradation rate and higher maximum degradation temperature) as shown in Table 1. When compared to montmorillonite or LDH, MWNT prove to be more efficient in terms of reduction of peak of RHR, as shown in Table 2. The reduction reaches 66% while the sample with clay shows a reduction of only 49%. In the residual char structure produced by the combustion, an integrated structure with surface
Table 1 Thermal degradation data of various EVA40 nanocomposites at 4 wt% filler loading (20C/ min under nitrogen) [35]
Sample
Tmax (C)
Maximum rate of degradation (%/C)
EVA40 EVA40–4MWNT EVA40–4CNF
453 460 456
2.24 1.37 2.39
164 Table 2 Cone calorimeter data of various EVA19 nanocomposites at 3 wt% filler loading (35 kW/m2) [36]
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PRHR (kW/m2)
Reduction (%)
EVA19 EVA19–3MWNT EVA19–3LDH EVA19–3MMT
1772 597 1090 903
– 66 39 49
± ± ± ±
170 30 58 24
cracks was formed from the clay/EVA nanocomposites while only limited char fragments were obtained in the composite containing the nanotubes. These structures do not appear to support the results on the peak of heat release rate if the fire retardant action of the nanotubes occurs in the condensed phase.
4.1.3 Combinations The use of the cited nanoadditives in elastomer nanocomposites proved to have a positive effect in terms of thermal resistance and flame retardancy. Nevertheless, it seems to be possible to combine the properties of the nanoparticles to obtain more resistant composites. That is why many researches are dedicated to potential synergistic effects between nanoparticles. A promising combination in terms of flame retardancy is the clay-carbon nanofiller mixture. Thermoplastic polyurethane elastomer was modified with different loadings of montmorillonite nanoclays and carbon nanofibers (CNFs) [37]. It was found that thermoplastic polyurethane elastomer with 10 wt% CNFs and with 5 wt% nanoclays gave the best thermal performance with respect to protecting a substrate. Indeed, the combination of the two nanoparticles resulted in the formation of a char layer, thus increasing the thermal insulative properties of the material. The same combination was reported in the literature in EVA [38, 39]. It appeared that the PRHR reduction is slightly increased when the two particles are nanodispersed in the matrix, but the aspect of the remaining char is radically different. The char produced by the clay-composite showed cracks on its surface, while the combination provides a smooth-surfaced carbonaceous shield. The roles of composites have been studied by Gao et al. It has been found that nanotubes may act as nucleation agents for graphitisation, leading to the formation of turbostratic and graphitic carbons. The formation of graphitic carbon in char may contribute directly to the reduction of the peak of heat release rate of the composites. This effect has been enhanced when both nanotubes and clay are used, since clay enhanced nanocomposites appear to have better resistance to char oxidation. The nanotubes also have the function to reduce surface cracks of chars, leading to the increase of barrier resistance to the evolution of flammable volatiles and to limit the oxygen transfer to the condensed phase. The use of carbon nanotubes with MMT in nanocomposites seems of interest as MWNT can provide mechanical properties to the char produced by the composite
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during combustion. That leads to an increase of the shield-playing role of the char and also of the fire resisting properties of the material.
4.2 Influence of the Dispersion The efficiency of the nanoadditives in polymeric materials is directly linked to their dispersion state. The presence of tactoïds or the total exfoliation of the particles influences the thermal or fire behaviour of nanocomposites, as reported by George et al. [35, 40] about nano-graphite dispersed in EVA60 (60% vinyl acetate). Transmission electron microscopy (TEM) pictures revealed that the natural graphite (NG) particles formed agglomerates inside the rubber matrix at 4phr as well as 8phr, on the contrary to expanded graphite (EG) platelets which were homogeneously distributed in the EVA matrix (Fig. 3). The given explanation was that the greater surface area of the expanded graphite can lead to higher interaction between the platelets and the macromolecules. These results were correlated to the measured thermal properties: the presence of NG adversely affected most of the properties of EVA while the incorporation of expanded graphite provided improvement in mechanical properties, thermal conductivity and thermal stability of EVA. A shift of 14C in maximum rate of degradation was observed for 4phr EG addition. The main parameter responsible for the dispersion quality in the composites is the compatibility of the nanofillers with the polymer matrix. So, to improve the dispersion of the nanoadditives it is important to increase this compatibility by modifying the filler or the matrix.
4.2.1 Compatibility with the Polymeric Matrix: Influence of the Surfactant The influence of the compatibility of EVA and different types of montmorillonite was investigated by Duquesne et al. [41]. They showed that the fire performance of
Fig. 3 TEM pictures of EVA60 containing a 4phr EG, b 8phr EG, c 4phr NG [Reproduced from Ref. 40]
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Fig. 4 SEM pictures of the EVA-Cloisite composites [Reproduced from Ref. 41]
EVA/C30B (Cloisite 30B is a MMT, organomodified by the incorporation of an alkyl ammonium ion) is very good at low loading (50% PRHR decrease at 5 wt% loading) whereas in the case of EVA/Na+ (pristine MMT), the improvement is lower (20% decrease for the same loading). In a second part, it has been demonstrated that the Cloisite Na+ is less compatible with the polymer because the ion is hydrophilic. On the contrary, the compatibility of Cloisite 30B in EVA is improved by the presence of the ammonium ions between the clay-layers. So, the structure of the composite obtained with Cloisite 30B is a nanocomposite-type structure whereas in the case of Cloisite Na+, a micro-structure is obtained. This assumption has been checked by small-angle X-ray diffraction analysis and by SEM analysis (Fig. 4). As a consequence, it was possible to conclude that a nano-structure enables to achieve better fire performance than a micro-structure. In fact, the presumed ‘‘diffusion effect’’ which leads to such an improvement, occurs in a nano-structure but not in a micro-composite. The work of Tang et al. [42] confirms the previous results. EVA/clay nanocomposites have been synthesized by melt intercalation from pristine MMT by adding compatibilizer C16 (hexadecyltrimethylammonium bromide). Figures 5 and 6 indicate that the clay compatibilizers have influence on the HRR. The initial Fig. 5 Heat Release Rate for pure EVA and different MMT composites (MMTa: nanocomposite, MMTb: microcomposite) [Reproduced from Ref. 42]
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Fig. 6 Heat Release Rate for pure EVA and EVA/MMT nanocomposites [Reproduced from Ref. 42]
HRR for the nanocomposite is higher at the beginning of combustion probably because of decomposition of the organic alky ammonium cation (C16) and of accelerated evolution of acetic acid owing to catalytic effect of acidic sites of the layered silicates. The modification of the surfactant of the clay layers is an efficient way to increase MMT dispersion and thus thermal and flame retardant properties. Another possibility to achieve better dispersion of additive having low compatibility with the elastomer is to modify the polymer matrix by grafting.
4.2.2 Use of a Compatibilizer: Grafting Since it does not include any polar group in its backbone, it is difficult to make intercalation of POE chains into clay layers. According to the literature, incorporation of polymer functionalized with maleic anhydride or hydroxyl groups as compatibilizer has been proven as a successful way to facilitate interactions between these two dissimilar components [9, 10]. This compatibilization was tested between maleic anhydride grafted POE (POE-g-MAH) and organo-modified montmorillonite to investigate the influence of MMT on the structure and flammability properties of POE-based nanocomposites [43]. The grafting allows a better dispersion of the MMT layers: the tactoïds present in the raw composite disappear in the grafted composite, where a complete exfoliation is achieved (Fig. 7). The enhancement of the dispersion quality is accompanied by an increase of the flame retardant properties. As shown in Table 3 the PRHR decrease of the grafted (and thus exfoliated) nanocomposite is higher than that of the composite presenting MMT tactoïds. The use of maleic anhydride as a compatibilizer for MMT was also reported in ethylene-propylene diene monomer EPDM [44]. Dispersion was characterized through X-ray diffraction (XRD) analyses, which showed an enhancement of the
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Fig. 7 TEM pictures of a POE/MMT and b POE-g-MAH/MMT nanocomposites [Reproduced from Ref. 43]
Table 3 Cone calorimeter data of POE, grafted POE and their nanocomposites (35 kW/m2) [43]
Sample
PRHR (kW/m2)
POE POE-g-MAH POE/MMT POE-g-MAH/MMT
1,727 1,650 1,280 793
platelets dispersion. Cone calorimeter was used to investigate the reaction to fire of the sample. It indicated that the PRHR decreases from 1374 kW/m2 for pure EPDM to 906 kW/m2 for the formulation with the compatibilizer whereas no PRHR decrease occurs with the organomodified clay. The compatibility of the nanoadditive in the polymeric matrix has a real impact on the thermal and fire properties of the resulting composite, since it directly influences the filler-polymer interactions and the dispersion quality. So, this compatibility is key parameter to achieve the best flame retardant performances. Thus, the synthesis of thermally stable or flame retardant elastomeric nanocomposites has to be designed to provide the best dispersion state, so as the nanoparticles can fully express their flame retardant qualities.
5 Nanoparticles in Elastomers Some nanoadditives are incorporated in polymeric matrices at relatively high loadings (typically 60 wt% for nano-MDH). However there is an interrogation concerning the relevance of incorporating nanoparticles at such loadings: making a calculation of an ideal dispersion (nanoparticles evenly dispersed individually) will reveal that the distance between 2 molecules is too low to avoid the agglomeration of particles. So the term ‘‘nano’’ is not suitable in that case, except for the elementary size of the particle.
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However, we propose to enlarge the notion of ‘‘nanocomposite’’ to composites filled with nanoparticles but where nanodispersion is not always achieved. That is why this part of the chapter is dedicated to the elastomeric matrices, flame retarded with nanoadditives but not nanodispersed.
5.1 Nanofillers in Elastomers Three major types of nanoadditives are used at high loadings in elastomers: the metal hydrates, the layered double hydroxides and calcium carbonate or sulphate.
5.1.1 Metal Hydrates and LDH The flame retardancy and mechanical properties of EVA containing 18% vinyl acetate (EVA18) containing Mg(OH)2 have been investigated by Huang et al. with reference to the particle size of MDH [45]. At a filling level of 55 wt%, there is little difference in mechanical properties and LOI values among four composites varying the particle size of MDH. The fire retarded blends, whatever the nano or micro-size MDH exhibit equivalent fire retardant properties. According to these results the expected effect of particle size does not exist. But it could be explained by the poor dispersion of nano-MDH observed in EVA compared with micro sized MDH since nano-size particles aggregate easier than micro-size one. At the opposite Lu et al. [46] with lamellar-like nano MDH noticed interesting properties. If UL-94 is not modified with nanoparticles compared to the same microparticles, the LOI is sharply increased (46 vol% for nano-MDH versus 39 vol% for micro-MDH at the same loading of 150phr). This might be explained by the better dispersion of the nanoparticles compared with microparticles, even if the presence of tactoids is detected by TEM analyses. It appears that dispersion is an essential parameter determining the efficiency of the flame retardant effect of MDH. Considering the need of mechanical properties and flame retardancy of composites, it is suggested that the smaller sized MDH should be selected as flame retardant, but the appropriate methods must be conducted to improve the dispersion of smaller particles. Concerning LDH, they have been tested in EVA with 14% vinylacetate (EVA14) [14] and proved to be efficient, as a LOI value of 42 vol% and a V-0 UL94 rating are reached with a filler content of 150phr (Table 4).
5.1.2 Calcium Carbonate and Sulphate Calcium carbonate mode of action implies the release of CO2 gases during combustion, which cool down the polymer substrate and in the same time dilute
170 Table 4 Effect of nano-LDH concentration LOI and UL-94 rating of the EVA14–LDH blends (specimens 117 9 12.7 9 3 mm) [14]
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UL-94 rating
Dripping
LOI (vol%)
100–50 100–80 100–100 100–120 100–150
Fail Fail Fail V1 V-0
Yes Yes No No No
27 32 37 40 42
flammable gases. Therefore the heat feedback to the polymer substrate is reduced and the emission of flammable gases decreases. Moreover, calcium carbonate contributes to the constitution of a mineral or ceramic-like calcium silicate residue after degradation which acts as flame retardant by a barrier effect [47]. Nevertheless this additive has to be incorporated at relatively high loadings to be efficient in polymeric matrices. Hamdani et al. [48] reported the use of calcium carbonate in a silicone matrix as a flame retardant. Thermo gravimetric analyses (TGA) typically show enhancement of thermal stability of silicone by addition of CaCO3 (30 wt%). Indeed, PDMS degrades around 300C, whereas when mixed with calcium carbonate, some residue remains in samples treated at 500C. The silicone elastomer stabilized by the additive thus resists at higher temperatures in addition to taking an active part in the formation of the protective structure. After treatment at 500C, the intumescent structure only consists of calcium carbonate and silicon oxides. This additive is also known to be used in styrene-butadiene rubber SBR. Nanosize-CaCO3-filled rubber showed a significant reduction in flammability in comparison with that without filler [49] according to the ASTM D4804 (the sample is clamped at a 45, a free end is exposed to a specified gas flame for 30 s, then the time required for the burning and the relative rate of burning are measured). The rate of flame retarding of nano-CaCO3 was more than that of commercial micrometric CaCO3. In this case, the rate of flame retarding for 9-nm CaCO3 was 2.55 s/mm, whereas for commercial CaCO3 it was 1.78 s/mm. The increase in flame retardancy with a decrease in the size was attributed to the homogeneous dispersion of the nanofillers. In the same matrix calcium sulphate CaSO4 at various loadings (between 2 and 12 wt%) and various particle size was compounded [50]. Thermal stability increased slightly when the particle size decreased: the temperature at the maximum degradation rate increases from 440C for neat SBR to 446C for filler loading of 4% (23 nm) or 460C at the same loading with a lower particle size (10 nm). The enhancement of the thermal stability is attributed to the homogeneous dispersion of nano-CaSO4 throughout the matrix (confirmed by scanning electron microscopy SEM). The nanoscale inorganic filler in the rubber nanocomposites also promotes the formation of the char layer, which acts as an excellent insulator and mass-transfer barrier, because of which this effect drastically improves the burning resistance.
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The particle size of the nanoadditives is also a crucial factor influencing the thermal stability or the reaction to fire of a material. Indeed, the specific surface increases with decrease particle diameter, leading to higher surface contact with the polymeric chains and thus to a modification of the composite behaviour. The effect of the particle size of CaSO4 was investigated in styrene-butadiene rubber [50]. The particle diameter was 10 nm, 15 nm, 23 nm or 20 lm. The thermal stability was evaluated through TGA measurements, which show an increase of the thermal stability with the decreasing particle size. The enhancement of the thermal stability is due to the uniform dispersion of nano-CaSO4 throughout the matrix, and the increased specific surface of the particle allow larger interactions with the polymeric chains, keeping them intact at higher temperature. The influence of the particle diameter also concerns the flammability of the material. The test ASTM D4804 reveals a lower flammability of the nanocomposites compared to the microcomposite, and this effect is more pronounced with low particle diameters. This is explained by the char-promoting action of the nanoscale inorganic additive in the rubber, which catalytic effect is higher with high specific area. The use of nano-MDH, LDH or nano-sulphate or carbonate in elastomers has been reported in the literature. These additives proved to be efficient in terms of LOI, mass loss calorimetry and UL-94 test, even if very high loadings (about 60% or more) are still needed to meet an adequate level of flame-retardant property. However, the nanoparticles are still detrimental to the mechanical properties. So, to develop their use in elastomers, it seems to be necessary to find synergistic systems allowing the decrease of the fillers content without decreasing the flame retardant properties.
5.2 Synergies and Use with Conventional Fame Retardant Additives Combining the nanofillers is an attractive way to reduce the overall fillers content in the formulation while achieving an acceptable level of performance, by the creation of synergies. As an example combinations of metal hydroxides with organomodified clay have been studied as a potential fire retardant system for polypropylene [51]. The combination of polypropylene with 20 wt% metal hydroxide and 5 wt% clay gives a reduction in the PRHR similar to that obtained with 40 wt% metal hydroxide. The use of nanofillers as synergists with other nanoadditives enables the creation of flame retardant elastomeric nanocomposites. Another way to enhance the thermal stability or flame resistance of such materials is the use of conventional flame retardant additives, whose action can be reinforced by appropriate nanosynergists.
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5.2.1 Combinations with Conventional Flame Retardants The incorporation of nanofillers to improve the flame retardancy of elastomers, flame retarded with conventional additives, is reported in the scientific literature [52]. In the case of elastomers the main combinations involve metal hydroxide and nanoadditives, such as clays, inorganic whiskers, carbon nanotubes and fumed silica. The decrease of the overall filler content is illustrated by the work of Beyer [53] about ATH, clays and carbon nanotube combinations. The RHR is maintained at 200 kW/m2 by the incorporation of only 5 wt% nanofiller while the ATH content is decreased from 65 to 45 wt%. Szep et al. [52] added modified (IMM = polybuthene/polysiloxane intercalated MMT) and unmodified montmorillonite (MMT) to EVA/MDH formulations. It appeared that the addition of MMT and IMM was detrimental to the UL-94 rating and slightly beneficial for LOI. The simultaneous use of modified and non-modified montmorillonite in EVA with magnesium hydroxide increases the effectiveness of flame retardancy (Table 5). The synergistic effect between MDH and MMT ? IMM was explained by the increased and sustained viscosity, by the early pre-carbonisation before the beginning of intense degradation and by the formation of a tough, stable ceramic-like residue from the montmorillonites and metal hydroxide interaction in the combustion process. Magnesium hydroxide sulphate hydrate whiskers/silicone rubber (SR) nanocomposites, were flame retarded with microencapsulated red phosphorus (MRP) used as a synergist [54]. It is found that whiskers can effectively improve the flame retardancy of SR composites due to endothermic degradation of whiskers with the release of water vapour diluting fuel supplied in the flame [55], accompanied by the formation of fibrous magnesia acting as a barrier layer. The incorporation of MRP in the system had a synergistic flame retardant effect in the condensed phase: the water released from whiskers promotes the formation of phosphoric acid from MRP, forming a protective phosphorus glassy layer [56]. A radical gas phase action is also suspected [57]. Silica particles can also be used as synergists in combination with metal hydroxides. The influence of fumed silica of different particle size on the fire retardant properties of EVA-ATH-MMT-silica blends was investigated by Laoutid et al. [58]. The partial substitution of ATH by 5 wt% of silica in an EVA/ATH/ MMT system did not improve significantly the overall fire behaviour. The presence of silica reduces the PRHR, especially for silica particles of small size and Table 5 LOI and UL-94 data of EVA28 compounds containing MMT and MMT–MDH [52] Material Weight ratio LOI UL-94 EVA ? MDH EVAMDH ? MMT EVA ? MDH ? IMM EVA ? MDH ? IMM ? MMT
1:1 8:8:1 8:8:1 16:16:1: 2
33 34 36 43
V-1 V-2 V-2 V-0
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large specific surface area, but reduces the appearance time of the second PRHR, due to the poor cohesion of the char formed with silica. The addition of silica in EVA-ATH-oMMT blends seems not beneficial regarding to fire resistance properties as it decreases char cohesion during combustion Flame-retardant methyl vinyl silicone rubber (MVMQ)/montmorillonite nanocomposites were prepared, using magnesium hydroxide (MH) and red phosphorus (RP) as synergistic flame-retardant additives, and aero silica (SiO2) as synergistic reinforcement filler [59]. In addition to LOI test, thermal properties were evaluated by thermogravimetric analysis (TGA). The decomposition temperature of the nanocomposite can be higher than that of MVMQ as basal polymer matrix. This kind of silicone rubber nanocomposite is considered by the authors a promising flame-retardant polymeric material.
5.2.2 Synergistic Combinations of Nanofillers Combinations with flame retardant additives are conventionally used to improve the flame retardancy of elastomers while decreasing flame retardant loading. In this part we will present the use of some exotic nanofiller which can be combined with other nanoadditives: the hydroxyl aluminium oxalate (HAO). HAO was tested as a synergist in clay/LDPE/EPDM composites [60]. NanoHAO (presented in Fig. 8) acts as heat remover and flammable gas diluter with the release of water and carbon dioxide after decomposing, while MMT mainly acts as heat and gas barrier, so the ternary composition has the aim to combine the two flame retarding properties, thus exhibiting superior flame resistance. It was found that the addition of MMT can produce synergistic effect on flameretarding nano-HAO/LDPE/EPDM system. The substitution of HAO by various amounts of clay leads to an increase of the LOI values. When the ratio of MMT and nano-HAO was 1:3, the LOI of the composites was up to 34 vol%, whereas HAO alone in the polymer matrix exhibits only a LOI of 31 vol%. Through the analysis of TGA, infra-red spectroscopy (FTIR) and SEM, it was proven that the barrier effect of the char layer was promoted when MMT was added and thus the decomposition of nano-HAO and polymer matrix under the char layer was retarded. The same approach was implemented with the combination of nano-kaolin and nano-hydroxyl aluminium oxalate in LDPE/EPDM composites [61]. From the LOI tests and UL94 tests, it was found that when 12 wt% nano-kaolin substituted nanoHAO in the composites, the LOI was enhanced from 31.0 to 35.5 vol% (Fig. 9), and the composites passed the UL94 V-0 standard, which proved that nano-kaolin
Fig. 8 Composition of nanoHAO
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Fig. 9 LOI values of the nano kaolin/HAO/LDPE/ EPDM composites (60 wt% filler content) [Reproduced from Ref. 61]
had a synergistic effect with nano-HAO on flame retardancy in the LDPE/EPDM system. Through the analysis of FTIR and SEM, it was found that the synergistic effect of nano-kaolin and nano-HAO on flame retarding LDPE/EPDM composites was attributed to the improved ability of rebuilding compact char barrier. Montmorillonite and nano-kaolin are efficient synergists in combination with nano-HAO. However, the very high nanoadditive loading (60 wt%) raises once again the question of the interest of using nanoparticles instead of microadditives, since the interest of nanofillers is to reduce the overall filler content without decreasing the flame retardant performances.
6 Conclusion In this chapter the design of thermally stable and/or flame retardant elastomeric materials was discussed. Two approaches were considered: the synthesis of elastomeric nanocomposites, where a real nano-dispersion is achieved, and the incorporation of nanofillers at high loadings, where agglomerates of nanoparticles are observed in the elastomeric matrix. Elastomer nanocomposites achieve a satisfactory dispersion state are flame retardant materials in specific conditions (fire testing involving mainly radiative phenomena) but generally fail flammability tests (LOI, UL-94). The overview of the nanoparticles mechanisms of action reveals that the formation of a protective layer during combustion is always involved whatever the used nano-additive. Moreover, the quality of the dispersion of the nanoparticles was found to be a key parameter in terms of flame retardancy. Since the dispersion is directly linked to the compatibility of the filler with the polymeric matrix, this aspect has to be considered, with the modification of the matrix (by grafting for example) or with the compatibilization of the filler (by coating or use of adapted surfactants).
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Moreover a particular attention should be paid to the mode of preparation of the nanocomposites, since it greatly influences the final dispersion state. The second approach suggested in this chapter concerns the incorporation of nanofillers in elastomers but at high loadings, thus preventing the nanodispersion. The use of nano-additives compared to micrometric ones was found more efficient in terms of flame retardancy or thermal stability, but some mechanical properties of the elastomers might be negatively influenced by such additives. The best flame retardant properties were obtained by the use of synergies between nanoparticles and combinations with conventional flame retardants. Promising developments are also expected in the creation of synergies, especially in nanocomposites. So, it is crucial to have a good understanding of the nanofillers mode of action, and of the type of synergisms created by the combination of nano-additives. Research work has to be focused on this particular technology as it provided nanocomposites exhibiting superior flame retardant properties.
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Recycling of Elastomeric Nanocomposites L. Reijnders
Abstract Recycling elastomeric nanocomposites is discussed in view of resource cascading, which aims at maximum exploitation of the value and service time of resources. Special attention is given to particulate releases linked to the elastomeric nanocomposite lifecycle. Recycling of tyres containing, at least partly nanoparticulate, carbon black is discussed as a practical example of the potential for resource cascading, recycling and particulate releases. In the case of elastomeric nanocomposites in general, prevention of degradation and exploring options for self-healing are worth considering in view of resource cascading. Recycling options for elastomeric nanocomposites include: reuse of the product, remanufacturing of the product and use of nanocomposite granulate (where appropriate, devulcanized). Further options are ‘chemical recycling’ (recycling of constituents or conversion products of nanocomposites) and incineration with energy recovery (‘thermal recycling’). Possibilities for the reduction of nanoparticulate releases linked to the elastomeric nanocomposite life cycle and socioeconomic arrangements favoring recycling are briefly outlined.
1 Introduction Firstly, I will clarify what ‘elastomeric nanocomposites’, also called nanocomposite elastomers, and ‘recycling’ actually mean in this chapter. Thereafter, in Sect. 1.3, the rest of this chapter will be briefly outlined.
L. Reijnders (&) IBED, University of Amsterdam, Nieuwe Achtergracht 166, 1018 WV, Amsterdam, Netherlands e-mail:
[email protected] V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_7, Ó Springer-Verlag Berlin Heidelberg 2011
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1.1 Elastomeric Nanocomposites Nanocomposite elastomers consist of elastomers and engineered nanoparticles, the latter being \100 nm in at least one dimension. In this chapter only inorganic nanoparticles will be considered. For instance, cellulose nanoparticles which are also considered for application in elastomeric nanocomposites [1] will not be dealt with, here. Use of engineered inorganic nanoparticles in elastomers has an already long tradition. Carbon black, traditionally used, and still the most popular filler, in rubber tyres, is at least partly nanoparticulate [2–5]. There is increasing interest in the use of inorganic nanoparticles in elastomers to confer improved performance regarding e.g. tensile strength, tear strength, abrasion resistance, crack growth resistance, barrier function, stiffness, hardness, permeability control, flame retardancy, electrical conduction, weatherability and stability [4, 6–31]. The nanoparticles studied in this context include ‘nanoclay’ (e.g. montmorillonite, fluorohectorite, bentonite, kaolinite), silica (SiO2), CaCO3, titania (TiO2) and carbon nanofibers and nanotubes, Silica particles are now commonly used in nanoparticulate elastomers [4]. In producing nanocomposites, often compatibilizing or functionalizing substances are used to promote dispersion of nanoparticles into polymeric matrices. Such compatibilizers include: (quaternary) alkylammonium compounds, alkylphosphonium compounds, polyolefin elastomers, maleated compounds and long chain amines.
1.2 Recycling Recycling is a loosely used concept. As such, it covers a range of activities varying from product reuse to incineration combined with energy recovery (which has been called ‘thermal recycling’). Recycling is important to the environmental performance of products (e.g. [32– 34]). To the extent that elastomer nanocomposites can be recycled a number of times as product, material or constituent substances with limited inputs into the recycling process, the environmental performance per cycle of usage might improve. The environmental benefit of many use cycles is at variance with a strategy aiming at biodegradable nanocomposite elastomers, which involves biodegradation after one cycle of usage [35–39]. Nanocomposites have drawn special attention in this context because nanoparticles may enhance biodegradability [37]. From the environmental point of view a strategy aiming at biodegradation would only seem worth considering when inputs in the recycling process are very high and the hazard of substances released by biodegradation, which in the case of nanocomposite elastomers would probably include hazardous nanoparticles (see Sect. 3), is very low. It would seem hard to come up with a case in which such conditions would be met by elastomeric nanocomposites.
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The applicability of recycling to elastomeric nanocomposites will be discussed against the background of a normative strategy regarding the resources used for producing nanocomposite elastomers. This normative strategy is based on the argument that recycling should aim at ‘resource cascading’, which in one version has been specified as the maximum exploitation of the value and service time of resources [40]. Alternatively, it has been argued that recycling should aim at maximization of the number of times that a material is utilized [41]. The difference between these two specifications is that the former values a more demanding application higher than a less demanding application and relatively long usage higher than short usage [42], whereas the latter gives equal value to all uses. Here the first specification will be followed.
1.3 Outline of the Chapter As pointed out in Sect. 1.1, tyres are a longstanding example of elastomeric nanocomposites. They belong to the category of thermoset elastomers, which one would expect to be more difficult to recycle at a high level of resource value than the category of thermoplastic elastomers [43]. Notwithstanding the latter, there is a long standing tradition of recycling tyres. So, to get a practical perspective on the recycling of nanocomposites, it would seem useful to look at current tyre recycling practices and the extent to which these are in conformity with resource cascading, as defined in Sect. 1.2. This will be done in Sect. 2. Also in Sect. 2 environmental effects of current usage and recycling of tyres will be outlined, including effects on the workplace environment. In this section special attention will be given to releases of particles. Section 3 will consider other nanoparticulate materials than carbon black which are currently used, or studied for use, in elastomeric nanocomposites. The nanoparticles which will be discussed in this section are: silica nanoparticles, carbon nanofibers and nanotubes, nanoclay and nano TiO2. This section will also consider what is known about the hazards of these nanoparticles to human health. Section 4 will consider the resource cascade and associated recycling options for elastomeric nanocomposites other than rubber tyres, and the potential to limit particulate hazards linked to these elastomers. Section 5 will present concluding remarks with a focus on the socio-economic arrangements which might be conducive to recycling.
2 Recycling of Tyres and Resource Cascading Tyres commonly contain a minimum of 20–30% carbon black and/or silica, which at least partially has a nanoparticulate character [4, 5, 8].
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Worldwide about 65% of all end-of-life tyres are landfilled or dumped, whereas *35% are the object of some form of recycling [44]. The main types of recycling are: retreading, recycling as rubber powder, rubber granulate or crumb rubber, and combustion with energy recovery [45]. Rubber powder or granulate, whether or not subjected to further treatment to improve performance, may find its way to a wide variety of applications, ranging from new tyres, brake pedals, rubber wheels, carpet tiles and livestock mats to athletic tracks and asphalt paving (e.g. [46–52]). The secondary product is often made of rubber. The properties of the secondary product are usually such that its performance tends to be at least somewhat below the standards required for the original tyre, even if there is a devulcanization step for secondary rubber and admixture of virgin rubber [53, 54]. To optimize product performance, the modulus of the recycled granulate should be matched to that of the virgin materials in the matrix [55]. Applications of rubber granulate may also be in composites of rubber and other materials as in the case of some varieties of asphalt and of rubber/polyolefin composites [56–59]. Renewed recycling of the latter type of composites should evade high temperatures because such temperature may degrade rubber and this in turn may negatively affect important properties such as tensile strength, tear strength and hardness [58]. ‘Thermal recycling’ is often in cement kilns, pulp and paper mills and power stations [49] (Beukering et al. 2001). Retreading and use of rubber powder or granulate tends to be favored by relatively high prices for mineral oil and new tyres [45]. Use of the major recycling options varies however widely among countries [45, 46, 60]. Of the main recycling options, retreading, which is in the more general category of remanufacturing [61], would come out relatively high in the resource cascade, directly after reuse of the product-as-it-is. Use of rubber powder or granulate would be lower in this resource cascade, and incineration with energy recovery still lower (cf. [62]). There are other options for recycling than retreading, use of rubber powder/ granulate and ‘thermal recycling’. The first thereof is pyrolysis. Pyrolyis of tyres has been studied extensively but has been applied commercially in a very limited way. Pyrolysis (also called thermolysis) of tyres may produce products varying from carbon black and oil to methane, hydrogen and monomeric feedstocks [44, 49, 63–68]. A problem that may arise in the pyrolysis of tyres is the presence of contaminants which interfere with further use and/or are relatively toxic. Examples of the former are N-compounds present in pyrolysis oil which can poison catalysts [69] and deposits on carbon black which necessitate further treatment before useful application [70]. Pyrolysis of tyres may also generate chars, which may be converted into activated carbon [71]. Dependent on the output of non-fuel products, pyrolysis may score higher in the resource cascade than incineration with energy recovery (cf. [62]).
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There are also other, in practice, minor recycling options. These include the construction of artificial reefs, the use in playground equipment, erosion control, bumpers for boats and highway barriers [46]. Finally there is research into self-healing of rubber [72] which is interesting in view of resource cascading. Self healing of tyres is as yet without substantial commercial applications. When such applications would come about, retreading would in the resource cascade perhaps come after self-healing tyres, depending on the actual composition of the tyre and the implications thereof for recycling (see Sect. 4.2). Important to resource cascading is that resources are conserved during a product life cycle. The tyre life cycle which includes recycling is associated with substantial resource losses, reflected in releases into the workplace and the wider environment [45, 73]. Such releases may have negative impacts on health. Also some of the recycling options, such as for instance the usage of tyre granulate in asphalt, may be associated with substantial resource losses. As to the health impact of resource loss, the following may be noted. Occupational exposure to carbon black has been found associated with increased pulmonary and cardiovascular illness, and exposure to carbon black has also been linked to oxidative stress and mutagenic/genotoxic effects [1, 5, 74–77]. Carbon black has been classified as possibly carcinogenic to human beings [76]. A tyre tends to lose *10% of its weight before it is disposed of [45]. Most of this material comes from the rubber tread and a substantial part of the particles which are released is in the respirable category (particulate matter with a diameter \10 lm or PM 10) [73]. Indeed, at busy city roads *5–7% of PM 10 may be tyre particles [73]. PM 10 is likely to contribute to increased risk of pulmonary and cardiovascular disease for people chronically exposed to particulate matter levels common along busy roads [73]. In scrap-tyre shredding facilities in Taiwan respirable particulate concentrations have been found which, on chronic exposure, would entail considerable risk of pulmonary and cardiovascular disease [73, 78]. These particles have moreover been found to be mutagenic, which may entail increased cancer risk on exposure [73, 78]. Also linked to the tyre life cycle, is the loss into the environment of Zn (zinc), which tends to be present in tyres at a level of about 1.5% [79, 80].
3 Nanoparticles for Use in Nanocomposite Elastomers and Their Hazards to Human Health The nanoparticles which will be discussed in this section are: silica nanoparticles, carbon nanofibers and nanotubes, nanoclay and nano-TiO2. As to the latter, also TiO2 nanoparticles coated with SiO2 will be considered. In this type of TiO2 nanoparticles the photocatalytic activity is suppressed. Regarding the hazards of these nanoparticles to human health, the following has been found. There is
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evidence that the nanoparticles considered here, when inhaled, are hazardous to lung tissue and, dependent on size, to the cardiovascular system [81–85]. Hazard to the cardiovascular system is probably linked to translocation from the lungs. Such translocation may also negatively affect other organs. The actual hazard following inhalation is dependent on number, diameter, surface area, structure, surface charge, surface energy, agglomeration and shape of the nanoparticles [81, 83, 86–96]. Ceteris paribus, carbon nanofibers and nanotubes which are longer than 10 lm would seem to be more hazardous than e.g. spherical nanoparticulate carbon black because macrophages which clean the lungs would seem to have much more difficulty in clearing such carbon fibers and tubes than carbon nanospheres [87, 97] and because carbon nanotubes inhaled by mice can reach the outer lining of the lungs and cause scarring [98]. Also, ceteris paribus, anatase TiO2 nanoparticles seem to be more hazardous than rutile TiO2 nanoparticles [99, 100]. There is furthermore suggestive evidence that nanoparticles can be translocated from the nasal area to the central nervous system via the olfactory nerve, thus posing a hazard of inhaled nanoparticles to the central nervous system including enhanced inflammation [83, 101–105]. There is some evidence that the nanoparticles considered here may be hazardous after ingestion [43, 106]. This hazard regards the intestines and, after translocation from the digestive tract, other organs. Uncoated TiO2 nanoparticles and carbon nanotubes, especially when residues of catalyst are present, may be a health hazard on dermal exposure [84, 107–109]. TiO2 nanoparticles have also been found to exhibit immunotoxicity [110]. A main molecular mechanism underlying the hazard of nanoparticles to human health may be oxidative damage caused by nanoparticles [81, 83, 85, 102, 107, 111–117]. However, this is probably not the only mechanism underlying hazard, as SiO2-coated TiO2 particles are linked to a relatively high hazard of lung inflammation, which is apparently not caused by oxidative damage [118]. Similarly, the necrotic effect of anatase TiO2 nanoparticles is unlikely to be explained on the basis of oxidation reactions only [99]. In evaluating the hazard of nanoparticles, one should be aware of the possibility that nanoparticles may absorb hazardous substances, and thus may cause exposure to combinations of substances [86, 93, 119].
4 Resource Cascading and Recycling of Nanocomposites Other Than Rubber Tyres Proper recycling is dependent on the availability of suitable technology and socio-economic arrangements conducive to recycling. With respect to the latter, tyres are not necessarily representative for all elastomeric nanocomposites. A first cause of differences is the matter of identification. Tyres are easily recognized and have a relatively well defined composition, and this may not hold for other
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nanocomposite elastomers. To remedy the problem of identification it has been proposed to tag products with information about composition which is relevant to recycling [120, 121]. Also in other respects there may be differences. For instance, car components may come under extended producer responsibility [45, 60], which is conducive to recycling, but this is not necessarily so for all nanocomposite elastomers. Designing optimum socio-economic arrangements for recycling is a complex matter (e.g. [122]). In this chapter the focus will be on the technical aspects of resource cascading and associated recycling. However in Sect. 5 the social and economic arrangements which might be conducive to recycling of nanocomposite elastomers will be briefly discussed.
4.1 Preventing Degradation of Elastomeric Nanocomposites Whether degradation during usage and recycling can be limited, is important to retaining high value for resource cascading in the case of nanocomposite elastomers. When oxygen is available, mineral nanoparticles such as SiO2, TiO2, C, and clay nanoparticles, generate reactive oxygen species, which may cause scission of polymer chains. An increased size of surface area, specific crystal structures and the chemical nature of the surface are important determinants of the capacity to generate reactive oxygen species [3, 5, 83, 112, 123, 124]. Polymer scission by reactive oxygen species may be suppressed by antioxidants. In the presence of sunlight, scission of polymers due to the formation of reactive oxygen species may also occur when nanoparticles turn out to be photocatalytically active. The latter applies to uncoated TiO2 nanoparticles [108] and to montmorillonite nanoclay [125–128]. Such nanoparticles may, when the nanocomposite is indeed exposed to sunlight (or another UV source), lead to a decrease in durability of the polymer [125–127], which in turn negatively affects the potential for resource cascading as defined in Sect. 1.2. The photocatalytic properties of the TiO2 nanoparticle can be suppressed by coating by e.g. silica, which for this purpose should preferably be complete [108]. Alternatively, increased additions of anti-oxidants, HALS light stabilizers and UV absorbents may be employed to reduce the deterioration of polymers by photocatalytically active nanoparticles. Degradation may also be linked to thermo-oxidative stress (e.g. during processing) and to delamination in the case that layered silicates (e.g. nanoclays) are used [124, 129–132]. The impact of mineral nanoparticles on the actual thermooxidative degradation of elastomers is somewhat uncertain as empirical studies have given rise to contradictory results [124, 132]. However, it seems that mineral nanoparticles are likely to be conducive to thermo-oxidative degradation [124, 129–134]. This may be linked to the generation of reactive oxygen species by mineral nanoparticles. It also appears that compatibilizers used in nanocomposites,
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such as long chain amines, quaternary ammonium salts and maleated compounds in the case of nanoclay, may enhance degradation, especially at elevated temperatures which might for instance occur during reprocessing [130, 132, 133, 135]. This has triggered a search for compatibilizers that have better thermal stability [135]. As degradation of nanocomposites may be linked to thermo-oxidative stress [132], there is a preference for low temperature recycling processes, which may also be preferable because such processes are associated with low energy inputs. Short residence time and limiting exposure to oxygen in processing elastomers during recycling may also be helpful in limiting thermo-oxidative stress [43]. Additives used to suppress oxidative damage due to the causes discussed here, may migrate out of the product, which may lead to poorer stability of the nanocomposite over time. Thus, minimization of migration, also at elevated temperatures that may for instance be necessary for reprocessing, should be important in the choice of additives [43]. But even so, it should be borne in mind that increasing chemical complexity of the product linked to the addition of additives may be unfavorable to further resource cascading [43, 126].
4.2 Self Healing of Elastomers To the extent that product degradation does occur, self-healing of nanoparticulate elastomers may be conducive to maintaining a high value in the resource cascade. Though there seems to be as yet no commercial application of selfhealing elastomeric nanocomposites, there is increasing research into its potential. Experimentally, self-healing of neat polymers has been demonstrated for rubber (based on hydrogen bonding), for polydimethylsiloxane (based on microcapsules with organotin catalysts) and for Diels–Alder elastomers (using encapsulated furan/maleimide) [72]. Also, there seems to be potential for selfhealing nanocomposites with carbon fibers and ferrite particles [72]. From the point of view of further resource cascading it is important that the inclusion of substances conferring self-healing properties does not interfere with later stages in the resource cascade.
4.3 Recycling Options for Elastomeric Nanocomposites From the point of view of resource cascading, the best recycling option would be reuse of the same product and recycling of production scrap into products. In analogy to rubber tyres and where appropriate, the second best option for recycling would seem remanufacturing of the product. In the case of nanocomposites exposed to wear and tear, this might for instance imply providing the product with a new surface layer. Thirdly, one can consider the recycling of nanocomposite elastomer granulate or powder, in view of minimizing the environmental burden preferably with low
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additional inputs such as energy, restabilizers and compatibilizers (cf. [62, 136, 137]). Among these options, reprocessing, whether or not combined with virgin material and—where appropriate—devulcanization, generating nanocomposite products meeting criteria that are not much inferior to the original product, is an interesting option [55, 137–144]. There is at least some scope for this approach. Thompson and Yeung [132] have shown that reprocessing of a nanoclay-thermoplastic olefin elastomer nanocomposite did induce some degradation in the nanocomposite, and that damage was reinforced by use of maleated compatibilizer, but that the rheological and mechanical properties of the recycled nanocomposite remained significantly better than those of the neat elastomer. Shim et al. [145] studied the option of devulcanization of silica-filled poly(dimethylsiloxane) and found this to be suitable to recycling, though revulcanization did restore original properties largely, but not completely. An alternative possibility is to include elastomer granulate in a matrix of elastomer with a different composition, using compatibilizer to achieve acceptable properties [146, 147]. It should be noted though that meeting performance requirements in the case of renewed recycling of elastomers composed of more than one type of polymer may be more difficult than in the case that one type of polymer is present [43]. Fourthly, one might think of recycling constituents or conversion products of the nanocomposite ‘chemical recycling’. For instance, dependent on the type of elastomer, one might consider chemical decomposition, to reclaim the original constituents of the elastomer. This has been demonstrated for flexible polyurethane, based on processes such as aminolysis, hydrolysis, alcoholysis and glycolysis, which regenerate constituent polyols [148–152]. Such processes require severe reaction conditions (which may include much elevated temperatures), which entail considerable inputs, and tend to generate substantial amounts of degraded by-products [153]. Also, it has been suggested to degrade polyurethanes by esters of phosphonic or phosphoric acid, and subsequently treat the reaction products with propyleneoxide [154]. This would generate fire retardants and other building blocks for the production of rigid polyurethane foams [154]. Again this is a type of ‘chemical recycling’ which requires considerable inputs. Inputs in polyurethane recycling processes and the generation of degraded by-products can be reduced by changes in polyurethane composition, which allow for decomposition at room temperature [153]. There are also options which focus on ‘chemical recycling’ of a small part of the nanocomposite. For instance, in the case that the nanoparticles applied are very valuable, as would currently apply to carbon nanotubes, one might consider reclaiming the nanoparticles from the nanocomposite e.g. by solvolysis of the elastomer, as also suggested for reclaiming longer carbon fibers from composites [155–157]. When the focus is on recycling a small part of the nanocomposite, one would expect a relatively poor environmental performance. Pyrolysis of nanocomposite elastomers is still another option in the context of ‘chemical recycling’, as it is in the case of rubber tyres (see Sect. 2) and in the case
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of neat elastomers (e.g. [62, 158–160]). Pyrolysis requires relatively high inputs of energy (e.g. [155]. One might aim at different outcomes, such as activated carbon, synthesis gas, oil and monomers. There is also a catalytic variant of pyrolysis, which might narrow the variety of substances generated [161]. The presence of mineral nanoparticles, such as clay, SiO2 and TiO2 nanoparticles, may considerably change the degradation patterns of polymers on pyrolysis [159, 162–168]. The presence of mineral nanoparticles in the nanocomposite might for instance change the composition of the gases produced, including the yield of monomers, or increase the generation of chars. Apart from pyrolysis, one might consider gasification and nanocomposite hydrocracking [169]. In resource cascading, incineration or ‘thermal recycling’ is usually the lowest stage of the cascade [62, 160]. It might however be that, in case of elastomeric nanocomposites, some types of ‘chemical recycling’ (e.g. focusing on recycling a small part of the nanocomposite or ‘chemical recycling’ involving large inputs) are more of an environmental burden than well managed ‘thermal recycling’.
4.4 Reduction of the Release of Hazardous Particles Associated with the Nanocomposite Life Cycle and Reduction of Nanoparticle Hazard In the production stage of nanocomposites, the release of hazardous particles can be reduced by in situ generation of nanoparticles in elastomers, such as described for nanosilica by Mark [170], Banyopadhyay et al. [7] and Ikeda et al. [12]. In the usage stage, robustness of the nanocomposite against degradation (cf. Sect. 4.1) and wear and tear [4] might reduce the release of nanoparticles. In recycling contained processing might reduce the release of nanoparticles [171, 172]. To the extent that nanoparticles are released during the life cycle of nanoparticulate elastomers, there may be scope for reduction of hazard by appropriate coatings for, and functionalization and purification of, nanoparticles. For instance it has been argued that the hazard of TiO2 nanoparticles can be reduced by coatings [173]. The hazard of amorphous SiO2 nanoparticles may be reduced by functionalization with amidosilanes [174]. The hazard of carbon nanotubes can probably be reduced by purification, which leads to the elimination of metalcatalyst residues, and the hazard of carbon nanotubes can be modulated by introducing functional groups [90, 175, 176].
5 Concluding Remarks It has appeared there is scope for resource cascade-type recycling of nanocomposites, and low nanoparticle releases along the product life cycle. Whether the
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technical options available will be used, strongly depends on social and economical arrangements. Both the examples of recycling of tyres and of the recycling of plastics (which does probably not exceed 10% of virgin plastics worldwide) illustrate this [43, 45]. Box 1 summarizes social and economic arrangements conducive to the recycling of nanocomposite elastomers [42, 43, 122].
Box 1. Social and Economic Arrangements Conducive to Cascade-type Recycling of Nanocomposite Elastomers – – – – – – – –
High prices of virgin nanocomposites; A tradition of (product) design for recycling; Labeling to identify the chemical nature of the nanocomposite; Extended producer responsibility of nanocomposites; Bans on incineration and landfilling; Well-organized recycling centers; Separate collection of end-of-life products; Well developed ‘reverse-logistics’ of end-of-life products.
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Part III
Applications
Elastomeric Nanocomposites for Tyre Applications Kaushik Pal, Samir K. Pal, Chapal K. Das and Jin Kuk Kim
Abstract In this study the epoxidized natural rubber (ENR) and organoclay (Cloisite 20A) composites were prepared by solution mixing process. The obtained nanocomposites were incorporated in natural rubber (NR) and styrene butadiene rubber (SBR) blends in presence of varying types of carbon black as reinforcing fillers. Morphology, curing characteristics, mechanical and thermal properties were characterized and analyzed. Also, the wear characteristics of the nanocomposites against Du-Pont and DIN abrader were determined and discussed. The morphology of the organoclay incorporated in ENR shows a highly intercalated structure. ISAF type of carbon black shows a significant effect on curing and mechanical properties by reacting at the interface between SBR and NR matrix. Blends containing ISAF N234 type of carbon black shows high abrasion resistant properties against Du-Pont and DIN abrader.
1 Introduction 1.1 General In automobile industry, the design of engine and other mechanical components receives prime importance, while the tyres are often overlooked. Most means of K. Pal (&) and J. K. Kim Polymer Engineering and Science, School of Nano and Advanced Materials, Gyeongsang National University, Jinju, Gyeongnam, 660-701, South Korea e-mail:
[email protected]@gmail.com S. K. Pal Mining Engineering Department, IIT, Kharagpur 721302, India C. K. Das Materials Science Centre, IIT, Kharagpur 721302, India
V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_8, Springer-Verlag Berlin Heidelberg 2011
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transport run by the use of tyres, thus, tyre plays an important and effective role in transportation. In automobiles, tyre is an inseparable assemblage of materials with very wide range of properties whose manufacture demands great precision. Tyre industries consume more than 60% of the rubber product; and the prime factors of consideration are safety and tread life. Thus, 30% of the mechanical work available from the fuel is dissipated in the tyres. A reduction of 15% in tyre rolling resistance may improve fuel consumption by 4.5% [1]. Throughout every industry, whatever type of machine is used, the concerns are the same: lowering operating costs and maintaining or improving the level of safety on site. In both areas, tyres are an integral factor. Tyres are one of the greatest consumable costs in surface mining and in underground trackless operations, representing as much as 20% of the operating costs of some machines [2]. Increasing the abrasion resistance of rubbers and rubber products is one of the problems of the rubber industry and is particularly important for the tyre industry. An analysis of the reasons, why tyres wear out?, from the data of tests on hundreds of thousands of series of production tyres shows that from 60 to 90% of tyres go out of use because of wear in the tread. On modern scales of tyre production, every 10% increase in useful life means a saving of significant costs. It is therefore important to the economy of the countries that the quality of tyres should be improved [3].
1.2 What is Tyre? A tyre is a composite, in other words an inseparable assembly of materials with very different properties, whose manufacture demands great precision. Tyres, or tires (in American and British English, respectively), are either pneumatic enclosures, or solid items (including rubber, metals and plastic composites). They are used to protect and enhance the effect of road wheels. Pneumatic tyres are used on many types of vehicles, from bicycles, motorcycles, cars, trucks, to earthmovers and aircrafts. Tyres enable vehicle performance by providing for traction, braking, steering, and load support. Tyres provide a flexible cushion between the vehicle and the road, which smoothes out shock, and provides comfort (Fig. 1). Fig. 1 Structure of tyre
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It is made up of the following semi-finished products: (1) (2) (3) (4) (5) (6) (7)
The inner liner, The casing ply, The lower bead area, Bead wires, Sidewalls, Bracing plies, Tread The fundamental tyre functions are:
(1) Proving load carrying capacity, (2) Provide cushioning and enveloping, (3) Transmit driving and breaking torque, (4) Producing tractive force, (5) Provide dimensional stability, (6) Resist abrasion, (7) Provide steering response, (8) Have low rolling resistance, (9) Provide minimum noise. (10) Permit minimum road vibration, (11) Be durable and safe. 1.2.1 Materials for Tyre The basic materials for the production of tyres are: (1) Base Rubber: NR, SBR, NBR, PBR, PUR, XNBR etc. (2) Fillers: carbon black, China clay (3) Additives: sulfur, peroxide, process oil (aromatic or aliphatic), wax, accelerator (CBS, MBT, MBTS, TMTD, DCBS, TBBS, DPG etc.), accelerator activator (ZnO, Stearic acid etc.), antioxidant and antiozonant (IPPD, HQ, TQ etc.), silica, nano fillers (clay, nanotuber, fibre etc.) (4) Wire: steel, brass, nylon, polyvinyl, polyamide, polyester cord etc.
1.3 Tyre Life and the Causes of Tyre Wear There are many factors that influence tyre life. Some of these are: 1. Cuts 2. Contamination 3. Dual tyre matching 4. Tyre rotation 5 Vehicle maintenance
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a. Misaligned wheels b. Wheel balancing c. Mechanical irregularities 6. Rolling resistance 7. Inflation pressure a. Over inflation/under loading b. Under inflation/over loading 8. 9. 10. 11. 12.
Heat Incorrect alignment Grip Load Abnormal tyre wear
1.3.1 Damage and Wear of Tyres Following are some reasons leading to damage and wear of tyres: 1. 2. 3. 4. 5. 6.
Tread detachment Air pockets Sidewall cuts and rupture of the sidewall Rupture resulting from a cut in the tread Impact ruptures Pinch rupture caused by road shock
Fig. 2 Different types of wounds in tyres
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7. Bead damage a. Incorrect fitting or removal b. Due to over inflation The images of several types of damages are depicted in the Fig. 2.
1.4 Phenomenon of Wear in Rubber Wear is ‘the process in which the tyre becomes unfit for use during a certain minimum service life’, which to considerable extent results from, or is related to, continuous damage of the tread [4]. Investigation by Rymuza [5, 6] shows that the wear dynamics of polymer– polymer and polymer-metal systems is determined by properties of the polymer such as surface energy, modulus of elasticity, specific heat, thermal conductivity and various operating conditions. Ratner [5, 7], Lewes [8], Rhee [9], Lancaster [10], Atkinson [11] and others have developed various forms of equations and relationships for the wear of polymers using variables such as load/pressure, speed, sliding length, sliding duration, shear strength of polymer etc. In 1974, Kar and Bahadur [12] developed a wear equation in terms of the sliding variables, pressure, speed, time and the material properties, modulus of elasticity, surface energy, thermal conductivity and specific heat. A proper understanding is necessary for control and prediction of polymer performance [13]. A vast amount of literature has appeared over the years in which relations between tribological performance and polymer properties are described in terms of mechanical parameters, such as yield and shear stress, toughness (as defined by the product of stress- and strain-to-break), plasticity index, Young’s modulus, and hardness [14–17]. These studies and reviews have helped in predicting polymer behaviour in sliding wear and friction to a certain level.
1.4.1 Relation between Abrasion Resistance and Mechanical Properties of Rubber Regarding physical ideas on the nature of abrasion, Schallamach [18, 19] was the first to examine the simple case of the failure of rubber by the action of a hard projection. The coefficient of friction l can also be expressed as a function of the elastic and hysteresis properties of the rubber and of the configuration of the abrading surface [18, 19].
1.4.2 Effect of Temperature on the Resistance to Wear In certain operating conditions of tyres e.g. sudden braking and acceleration, sharp bends and high speeds etc., high temperature develops in the contact area [20].
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It has been observed [19] that during intense abrasion in sliding contact, high temperature is developed, and consequently, the abrasion resistance of the rubber depends to large extent on the resistance to high temperature and heat. One suggestion is that chemical degradation of the sliding interface of the rubber can be attributed to a thermal effect [21]. In the wear of tyres, the temperature of the rubber at the interface needs to be considered in relation to abrability, which is a function of temperature [22]. 1.4.3 Mechanism of Wear of Tread Rubbers Tread wear in pneumatic tyres may be measured as weight loss or decrease in crown thickness, or more commonly as a loss of tread depth over prolonged periods. Pavement texture, as might be supposed, plays a major role in determining the extent and severity of the wear mechanism, in addition to driving habits, climatic conditions and operational factors. There are several mechanisms involved for tyre abrasion, such as, (1) (2) (3) (4) (5) (6)
Fatigue, or hysteresis wear Abrasive, or catastrophic wear Cohesive tearing, or wear by roll formation Tread reversion and blistering Smearing of rubber Threshold strength of rubber
1.4.4 Relation between Abrasion and Tensile Strength of Rubber The investigation of abrasion by scratching rubber with needle led Schallamach to the conclusion that wear of rubber on sharp abrasives was due to tensile failure [23, 24]. Buist and Davies [25] proposed an empirical relation between volume (V) of wear and physical properties of rubber as V ¼ C0 þ C1 Shore hardness þ C2 Tensile strength
ð1Þ
where, C0, C1, and C2 are constants. Thornly [26] has similarly been able to correlate tyre wear with hardness and tensile properties. The effects of the particle size and structure of various carbon blacks on friction and abrasion behavior of filled natural rubber (NR), styrene-butadiene rubber (SBR) and polybutadiene rubber (BR) compounds were investigated [27] using a modified blade abrader. Characteristic parameters like particle size and the structure of carbon blacks were observed to have a linear relationship with the Young’s modulus. The frictional coefficient depended not only on the particle size, but also on the structure of carbon black. The rates of abrasion were decreased with increasing surface area and developing structure of carbon blacks. Parkins [28] showed that abrasion resistance of the carbon black filled rubber would increase if the tensile strength is higher.
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Rattanasoma et al. [29] showed that the vulcanizates containing 20 and 30 phr of silica in hybrid filler exhibit better overall mechanical properties as well as abrasion resistance properties. Persson [30] reported that friction force decreased as the velocity increased in the high velocity range depending on the nature of the substrate, surface roughness and the mechanical properties of rubber. Persson’s attention was focused on examining the behavior of bulk rubber undergoing continuous motion. Many attempts have been made to find a relationship between the rate of abrasion loss and the physical and mechanical properties of rubber [22, 31–33]. Uchiyama [33] found the following relationship for determining the wear volume of rubber abrasion: V ¼ k1
lP L rB
ð2Þ
where V is the wear volume, l represents the friction coefficient, P denotes the normal load, L the length of rubbing distance and k1 is a constant. The parameter rB is expressed by the following equation: rB ¼ rN
ð3Þ
where r is maximum amplitude of tensile stress and N is the number of cycles. It has been shown that the wear resistance could be correlated with the mechanical properties of the vulcanizates [34].
1.5 Development of Wear Resistant Rubber Blends Polymer blends are being used extensively in numerous applications, especially in tyre manufacture. In the time of the twelfth century, the source of the whole of the rubber supplied to the industry was natural rubber. In the nineteenth century, Charles Goodyear invented the process of ‘sulfur vulcanization’ and suggested the use of ground natural rubber to overcome its limitations [35]. As a consequence of the first world war, Germany introduced Buna rubbers which are purely synthetic rubber and it increased the curiosity of polymer chemists all over the world. Various attempts were made in laboratory to enhance the properties of natural rubber thus transforming it into a material of desirable properties. As a result, synthetic rubbers with tailor-made properties were born. Consequently, different chemicals and methods for vulcanization and processing were developed. Apart from blends of common rubbers, specialty rubber is also utilized for such purposes, depending on service demands and components of the tyre [36, 37]. Many reports covering a wide range of rubber blends have been published. The use of carbon black is synonymous with the history of tyres. Although it
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has lost some ground to other reinforcing fillers such as silica, but, by virtue of its unrivalled performance, it is still the most popular and widely used reinforcing filler. However, the primary properties of carbon blacks are normally controlled by particle size, surface area, structure, surface activity and they are in most cases interrelated [38]. In tyre treads, silica can yield a lower rolling resistance at equal wear resistance and wet grip than carbon black [39]. Natural rubber (NR) is a virgin rubber, having properties resembling those of synthetic rubbers. Natural rubber (NR) is known to exhibit numerous outstanding properties; reinforcing fillers are necessarily added into NR in most cases in order to gain the appropriate properties for specific applications. A wide variety of particulate fillers are used in the rubber industry for various purposes, of which the most important are reinforcement, reduction in material costs and improvements in processing [40]. Reinforcement is primarily the enhancement of strength and strength-related properties, abrasion resistance, hardness and modulus. It can offer unique properties such as good oil resistance, low gas permeability, improved wet grip and rolling resistance, coupled with high strength. A lot of research has been carried out on NR and SBR blends by varying the quantity and composition of additives and fillers. The abrasion resistance of styrene-butadiene tread rubbers is observed to depend to large extent on the molecular weight distribution, the average molecular weight and certain other factors [41–43]. Solution styrene butadiene rubber (S-SBR) is used in a wide variety of applications, including the production of tyres, footwear, conveyor belts, hoses, flooring and adhesives [44]. Solid solution polymerized styrene butadiene rubber, produced by anionic batch polymerization is available in a wide variety of styrene and vinyl contents. S-SBR rubbers provide excellent balance between wet grip, rolling resistance and dry handling in silica and carbon black compounds for high-performance tyres [45]. They are also used for the manufacture of high quality technical rubber goods. As it is well known, the performance of motor car tyres must be improved, whilst at the same time reducing the amount of natural rubber used. This problem applies even more to large truck tyres, and can be solved provided that new types of stereo-regular synthetic rubbers like isoprene [46–48] and butadiene [46, 49–51] are used. The effects of mixing method, blend ratio, content, type of carbon black, and vulcanization system have also been compared. Tyres used in mining vehicles are very costly and need regular maintenance, and it is expensive to replace them within a very short term. The rugged working conditions in mining industries reduce the life span of tyres on account of cuts, contamination, abrasion, wear, speed fluctuations etc. There are several types of damage which occur in the dumptruck tyre such as tread detachment, sidewall cuts, impact ruptures, bead damage etc. [39, 52]. The idea of blending synthetic rubbers with natural rubber is certainly not a new one, but it is only now that this can be shown to be possible with consistently positive results, by the use of new techniques developed over the last years.
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1.5.1 Compatibility of Polymer Blends Polymer blend is the intimate mixture of two or more polymers and/or copolymers resulting from common processing steps [53]. Throughout the last decades, scientific and technical literatures in this area have expanded remarkably as evidenced by appearance of several books and proceedings of various conferences [54]. The word compatibility has been used by many investigators to describe single phase behavior. However, the terms compatible and incompatible refer to the degree of intimacy of the blends, which depends on the measurement procedure during examination. A blend could be considered as a compatible blend, if it does not exhibit the gross symptoms of component polymer segregation, whereas a heterogeneous blend at a macroscopic level is incompatible. Compatible system of blends containing high molecular weight polymers have been identified usually when a favorable specific interaction such as hydrogen bonding, dipole interaction or ionic interaction exists between two components. Although the majority of thermoplastics/elastomer blends are heterogeneous, the components may be referred to as compatible if some technically advantageous combination of properties can be realized from the blends. Partial compatibility implies that above a particular level either the minor or the major components remain as a dispersed phase. There are some technical problems, which are frequently the result of some type of mutual incompatibility, which provide an inferior set of properties when dissimilar polymers are blended together [55]. The blending process and the quality of the blends can be improved by adjusting the individual raw polymer viscosity. As a result, the effective viscosities of the phase will no longer mismatch. The thermodynamic incompatibility can be overcome if the surface energy differences between polymers are small enough to permit the formation of very small micro domains of the individual polymer phases and there is sufficient adhesion between the phases by formation of crosslinks across the interface during blending [56, 57]. The compatibility of various components and the generation of single phase from multiphase system play a major role in influencing the physical properties of the polymer blends. The most significant need of the designer of polymer blends is to ensure good stress transfer between the components of the multi-component system which can only vouch for the efficient utilization of component physical properties of the blends. Numerous techniques have been utilized to determine the compatibility of the blends but only a few predict good results [58, 59].
1.6 Nanofillers Nanofillers have for many years high significance in the plastics industry. Nanofillers are basically understood to be additives in solid form, which differ from the polymer matrix in terms of their composition and structure. They
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Fig. 3 Nanosilica as observed from TEM
generally comprise inorganic materials, more rarely organic materials. Inactive fillers or extenders raise the quantity and lower the prices, while active fillers bring about targeted improvements in certain mechanical or physical properties. The activity of active fillers may have a variety of causes, such as the formation of a chemical bond (e.g., cross linking by carbon black in elastomers) or filling of a certain volume and disruption of the conformational position of a polymer matrix, and also the immobilization of adjacent molecule groups and possible orientation of the polymer material [60]. There are many grades of nano fillers, e.g. carbon black, carbon nanotubes, carbon fiber, activated clay, natural clay (mined, refined, and treated), clay (synthetic), natural fiber, silica, zinc oxide etc. (Fig. 3). 1.6.1 Effects of Nanofillers Nano filler are expected to improve the properties of materials significantly, more even at lower loading than conventional/micro-fillers. There are some reasons for that [61]: (1) The size effect: large number of particles with smaller inter-particle distance and high specific surface area results in larger interfacial area with the matrix. (2) The interactivity and potential reactivity of the nanofillers with the medium.
1.6.2 Nanofiller Reinforcement The difference between the behaviors of micro and nano-reinforced polymers can be analyzed by observing the specific changes in properties in nanoscale, in which polymer chain lengths approach the filler dimensions so that they might display particular interaction influencing the macroscopic behaviour of the materials.
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Fig. 4 Different phases of nanocomposites
Indeed, many parameters can be taken into account for the reinforcement efficiency of filler into given medium (Fig. 4) [62, 63]. (1) (2) (3) (4) (5) (6) (7)
chemical nature of fillers shape and orientation of the fillers average size, size distribution, specific area of the particle volume fraction dispersion state interfacial area respective mechanical properties of each phase
1.6.3 Several Uses of Nanofillers • Solar energy—tougher, more efficient solar cells are already under development, with the promise of drastic cost reductions on the horizon. Some will even produce hydrogen. • Fuel cells • Display technologies and e-paper—e-paper and carbon-nanotube-based fieldemission displays expected to be slugging it out with liquid–crystal displays (with carbon-nanotube-based backlights, of course) in the next 2 years. • Nanotubes—both as raw materials and as products. Multi-walled nanotubes are already used in composites, to increase conductivity at much lower filler loads. Single-walled nanotubes will have a much bigger effect in the longer term. • Catalysis has a huge potential geopolitical impact, especially after recent developments in the energy business. • Nanocomposites—mainly clay-based for structural applications (increased strength) or with novel properties. These are already penetrating the automotive and aerospace industries.
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• Storage technologies—magnetic random access memory (RAM), nanotube RAM and terabyte hard drives in the next few years. • Nanocrystalline bulk materials or steels containing nanoparticulates—some companies are already using steel with nanoparticulate carbon added during the rolling process. • Coatings—extra hard or with special properties, such as being electrochromic or self-cleaning, are under investigation by everyone from car manufacturers to architects. • Sensors—bio and chemical sensors made from nanowires and nanotubes are currently probed. • Bio analysis—devices using atomic force microscopes and quantum dots are already being readied for market. • Textiles—nanofibres in stain-resistant trousers are already available, with electrospun nanofibres and nanotube-enhanced fibers coming soon.
1.7 Tyre Retreading Tyres that are fully worn can be re-manufactured to replace the worn tread. Retreading is the process of buffing away the worn tread and applying a new tread. Retreading is economical for truck tyres because the cost of the new tread is small compared to the cost of the tyre carcass. Retreading is less economical for passenger tyres because the cost is high compared to the cost of a new tyre. However, commercial truck drivers run the risk of ‘‘blow-outs’’, separation, and tread peeling from the casing, due to constant re-use of the casing. Commercial trucking companies have taken their own initiative as well. Many only run retreads on their trailers, and keep ‘‘Virgin Casings’’ (new tyres) on their Steer and Drive wheels. This ensures that in the event that a retread blows out, the driver maintains control over the truck.
1.8 Objectives and Scope of Work As mentioned earlier, apart from the blends of common elastomers, specialty elastomers are also utilized for tyre applications, depending on service demands and components of the tyre [36, 37] and carbon black is one the most commonly used fillers [44, 64–66]. The primary aspect in preparing organoclay nanocomposites is to attain a very high degree of dispersion of organoclay aggregates that afford to very large surface areas. Hence, it exhibits significant improvements in physical, mechanical and thermal properties in relation to the polymer host [67–69]. Though many organoclay thermoplastics have been prepared and studied [70–73], less attention has been
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paid to use organically modified layered silicates in reinforcing elastomers [74–76]. In this study, rubber-organoclay nanocomposites have been prepared and analyzed. The accomplishment of highly dispersed organoclay nanocomposites has two requirements. The first requirement involves the compatibility between the polymer and nanoclay to obtain better dispersion of organoclay in the polymer matrix. Organoclay can be easily dispersed in polar polymers than in non-polar polymers. The formulation of organoclay/polar polymeric systems usually contains a polymeric compatibilizer [77, 78] e.g. ENR. ENR is obtained by epoxidation of 1, 4-polyisoprene, depicts higher glass transition temperature and increased polarity. Accordingly, as a fine dispersion of organoclay is needed, ENR was also chosen as a compatibilizer in this study. Arroyo et al. [52], Teh et al. [79] and Varghese et al. [80] have carried out few resplendent works using ENR as compatibilizer for organoclay/natural rubber nanocomposites. In our previous literature, we have already analyzed the effect of ENR as a compatibilizer in natural rubber-nanoclay gum compounds [81] and in presence of carbon black [82]. The second requirement is the preparation methods of nanocomposites. Various methods have been adopted for the preparation of rubber/organoclay nanocomposites that include in situ polymerization intercalation [83], solution intercalation [84] and melt intercalation [85], and finally co-coagulation of rubber latex and clay aqueous suspension [86]. In this study, the solution intercalation process has been used to make the organoclay nanocomposites. Tyres used in mining vehicles are very costly and need higher abrasion resistance. An increase in the abrasion resistance of rubber products can be achieved by studying the mechanism of wear of rubber under different operating conditions. The wear of rubber is a complex phenomenon and dependent on a combination of processes such as mechanical, mechano-chemical and thermo-chemical etc. Schallamach [87] and later Grosch [88] reviewed abrasion of rubber and tyre wear. Champ et al. [89] and Thomas [90] suggested that abrasion takes place through a cyclic process of cumulative growth of cracks and tearing. Kragelskii et al. [91] and Schallamach [92] examined the simple case of the failure of rubber by the action of a hard projection moving over its surface. It has been observed that during intense abrasion in sliding contact, a high temperature is developed, and consequently the abrasion resistance of the rubber depends, to a large extent, on its resistance to high temperature and heat [92]. The earlier literature review has demonstrated several types of damages and its causes in the tyre [93–112]. Blending of elastomers has been often used to obtain an optimum number of desirable combinations, physical properties, processability and cost. The elastomers selected in this study were natural rubber (NR) and styrene butadiene rubber (SBR). Since NR and SBR are non-polar, epoxidized natural rubber (ENR) is used as a compatibilizer to improve the dispersion of organoclay for nanocomposite preparation. In this study, incorporation of organoclay in ENR was done by solution mixing, in order to obtain uniform dispersion of the nanoclay in ENR. The obtained ENR-organoclay nanocomposites were incorporated in the NR/SBR blends with varying types of carbon black. The cure characteristics, morphological, mechanical and thermal properties of nanocomposites were investigated.
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These rubber compounds were examined in a specially fabricated experimental set-up for evaluating their wear resistance properties, when abraded against various rock types.
1.9 Experimental 1.9.1 Materials Used in Rubber Preparation Natural rubber (RMA-1X) was supplied by the Rubber Board, Kottayam, Kerala. The styrene butadiene rubber (SBR) used was Krylene HS 260, No.-5 of 1948 grade of Bayer AG. Its styrene content is 23.5 ± 1.0, specific gravity is 0.94 and Mooney viscosity at 100C is 50 ± 5. Epoxidized natural rubber containing 47% epoxidation unit was supplied by Agricultural Product Processing Research Institute, Zhangiang, PR China. Cloisite 20A, a natural montmorillonite (MMT) modified with a quaternary ammonium salt with cation exchange capacity of 95 meq/100 g clay (Southern Clay, Inc, USA), was used as a nanofiller in the preparation of the nanocomposites. Zinc oxide, stearic acid, N-cyclohexyl-2benzothiazyl sulfonamide (CBS) and N-isopropyl-N-phenyl-p-phenylenediamine (IPPD) were supplied by Bayer (India) Ltd. Standard rubber grade process oil (Elasto 710) and paraffinic wax were purchased locally. Carbon black was supplied by Birla Carbon. 1.9.2 Solution Mixing Method Epoxidized natural rubber was dissolved in methyl ethyl ketone (MEK). The ratio of the rubber to solvent was 1:3 (weight/volume). Continuous stirring was performed at room temperature, until the rubber was completely dissolved in the solvent. Subsequently 100 wt% of nanoclay (Cloisite 20A) was added to the rubber solution and stirring was continued. The resultant solution was then cast over in a thoroughly cleaned plane glass plate. The sample was kept in the same condition until the solvent was completely evaporated. Appearance of a transparent film was observed. The obtained nanocomposites contained 1:1 ratio of ENR and nanoclay. 1.9.3 Preparation of Nanocomposites The compounds were prepared in two-roll mixing mill operated at room temperature. The speed ratio of the rotors was 1:1.4. Initially the natural rubber and styrene butadiene rubber was masticated followed by incorporating ENR/nanoclay composites. The reinforcing filler (carbon black) was added along with the process oil followed by curatives and shown in Table 1. Also, three types of carbon blacks and one semi reinforcing filler were used, such as, SAF N110, ISAF N231, ISAF N234 and SRF N774. For vulcanization, the
Elastomeric Nanocomposites for Tyre Applications Table 1 Compound formulation Compounds
Natural rubber (NR) Styrene butadiene rubber (SBR) Carbon blacks 1. SAF (N110) 2. SRF (N774) 3. ISAF (N234) 4. ISAF (N231) ENR/nanoclay (Cloisite 20A) Stearic acid Antioxidant (HQ) Accelerator (CBS) Zinc oxide Process oil Sulfur
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Sample codes A B Weight in wt%
C
D
E
F
75 25
75 25
75 25
65 35
65 35
65 35
20 20 – – 10 2 1 1 5 2 2
– – 40 – 10 2 1 1 5 2 2
– – – 40 10 2 1 1 5 2 2
20 20 – – 10 2 1 1 5 2 2
– – 40 – 10 2 1 1 5 2 2
– – – 40 10 2 1 1 5 2 2
amounts of additives such as sulfur, process oil, CBS were based on 100 wt% of rubber and the samples had the codes ‘A’, ‘B’, ‘C’, ‘D’, ‘E’, and ‘F’, respectively. The physical properties of the different types of carbon black used in this study have already been discussed in our earlier literature [113].
1.9.4 Experimental Techniques Experimental techniques followed in the present research are as follows.
Cure Characteristics of Rubber Compound The cure characteristics of the rubber compound were studied with the help of a Monsanto Oscillating Disc Rheometer (ODR—100 s) at 150C as per ASTM D-2084-07. From the graphs, the optimum cure time, scorch time and Cure Rate Index (CRI) could be determined.
Determination of Crosslink Density The cross-link density was determined by immersing a small amount of sample in 100 ml benzene for 72 h to attain equilibrium swelling. After swelling, the sample was taken out from benzene and the solvent was blotted from the surface of the sample and weighed immediately. This sample was then dried out at 80C to remove all the solvent, and reweighed. The volume fracture of rubber in swollen
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gel Vr, which represented the relative cross-link density of the vulcanizate, was determined by the following equation. Vr ¼
m0 Uð1 aÞ=qr m0 Uð1 aÞ=qr þ ðm1 m2 Þ=qs
ð4Þ
where m0 is the sample mass before swelling, m1 and m2 are sample masses before and after drying, U is the mass fraction of the rubber in the vulcanizate, a is the loss of gum EVM vulcanizate during swelling, qr and qs are the rubber and solvent densities respectively. The physical cross-link density was calculated using modified [114] FloryRehner equation from swelling measurement in benzene reported earlier [115]. 1 ½Vr þ vVr2 þ lnð1 Vr Þ ¼ l 1=3 Mc qr VS ðVr Vr =2Þ
ð5Þ
where, 1/M’c, Vr, Vs, v and qr are the physical cross-link density, volume fraction of rubber, molar volume of swelling medium, the Flory–Huggins solvent-rubber interaction parameter and density of rubber respectively. Lastly, the chemical cross-link density was determined from the following equation. 1 1 1 ¼ þ 1:55 105 l 2Mc 2Mc M
ð6Þ
where 1/2Mc, 1/M’c and M are the chemical cross-link density, physical cross-link density and molecular weight of rubber vulcanizate respectively.
X-Ray Diffraction Measurements (XRD) X-ray diffraction was performed with a PW 1840 X-ray diffractometer with a copper target (Cu-Ka) at a scanning rate of 0.050 2h/s, chart speed 10 mm/2h, range 5,000 c/s, and a slit of 0.2 mm, applying 40 kV, 20 mA to assess the change of crystallinity of the blends as a function of blend ratio [116]. The range of 2h scanning of X-ray intensity employed was 1.5–10. The degree of crystallinity (vc) was measured using the following relationship: vc ¼ Ic =ðIa þ Ic Þ
ð7Þ
where, Ia and Ic are the integrated intensity of the crystalline and amorphous region, respectively. The crystallite sizes (P) and the interplaner distance (d) are calculated as follows: P ¼ Kk=b cos h
ð8Þ
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ð9Þ
where b is the half height width (in radian) of the crystalline peak, k is the wavelength of the X-ray radiation (1.548 for Cu), and k is the Scehrrer constant taken as 0.9 [117]. Mechanical Characterization (Tensile and Tear) Vulcanized slabs were prepared by compression molding, and the dumb-bell shaped specimens for tensile tests and crescent shaped specimens for tear tests were punched out from a molded sheet by using ASTM Die C. The tests were performed on a universal tensile testing machine (Hounsfield H10KS) under ambient condition (25 ± 2C), following the ASTM D 412-06 and ASTM D 62400 (2007). The modulus at 100, 300 and 500% elongation, tensile strength, tear strength and elongation at break (%) were measured at room temperature. The initial length of the specimens was 25 mm and the speed of the jaw separation was 500 mm/min. Samples were tested five times for each set of conditions, at the same elongation rate. The values of the tensile strength, modulus at 100% elongation, 300% elongation, 500% elongation and elongation at break were averaged. The relative error was below 5%. The hardness was measured by Shore A hardness tester following ASTM D2240-05 standards. Thermal Characterization Thermal characterization (TGA) studies were carried out Shimadzu-DT-40 instrument in presence of air at a rate of 10C/min, using temperature range of 25 to 650C. Degradation temperature of the composites was studied through this analysis. It is well known that because of the high flexing of tyre off the road (dumptruck), the temperature rises. In tropical countries, the temperature can go as high as 150C. Hence, the study of thermal resistance of the compounds is desirable. Transmission Electron Microscopy (TEM) The dispersion morphology was observed in the high resolution transmission electron microscope (HRTEM, JEOL 2100). The samples were ultramicrotomed at -20C for ENR/nanoclay films and -70C for ENR/nanoclay composites in NR.
Scanning Electron Microscopy The tensile fracture surface of the samples was studied in a scanning electron microscope (JSM-5800 of JEOL Co.; acceleration voltage of 20 kV; gold coating
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and 500 and 1000 times magnification. Scanning electron microscopy (SEM) was used to study the morphology like filler dispersion and indentation of the abrader on the abraded rubber surface of the samples prepared.
Du-Pont Abrasion Test Du-Pont abrasion test was done by the Du-Pont Craydon type of abrasion tester for determining the abrasion resistance of compounds of vulcanized rubber recommended by the Indian Standards Institution vide IS:3400 (Part-III)—1965.
DIN Abrasion Test DIN abrasion test was done by the DIN abrasion tester for determining the abrasion resistance of compounds of vulcanized rubber recommended by the Indian Standards Institution vide IS:3400 (Part 3)—1987.
Heat Buildup Study Heat build up study was carried out using Goodrich Flexometer for the selective compounds having higher tensile strength. The test pieces were prepared in cylindrical shape having diameter 17.8 ± 0.15 mm and height of 25 ± 0.25 mm by compression molding at 150C. The test pieces were kept at initial temperature of 50C and stoke of 4.45 ± 0.03 mm. The temperature and the load were kept constant throughout the study. The temperature attained by the samples after the time periods of 10 and 20 min was recorded.
1.10 Results and Discussions 1.10.1 Cure Characteristics of the Rubber Compounds The optimum cure time (t90) for ‘C’ and ‘E’ rubber sample is higher than other rubber vulcanizates as shown Table 2. It is possibly due to the mixing of 25 wt% of SBR irrespective of carbon black grade. The t90 of the compounds sharply decreased may be due to the amine functionalities in the filler after the modification process or ion exchange process. The rate of cure {tmax - tmin} always increased with increasing concentration of NR. This increase in cure rate can be due to the fact that an increasing concentration of NR caused the increase in vulcanization reaction and created more active cross link sites in the rubber compound.
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Table 2 Cure characteristics Properties
A
B
C
D
E
F
Min. Torque (dN m) Max. Torque (dN m) TS2 induct time (min) TS5 scorch time (min) TC90 opt. cure time (min) Opt. cure (min) Cure rate index (min-1) Crosslink density (moles/g) 9 10-5
11.39 37.58 1.71 2.32 5.64 35.23 23.86 2.27
12.19 32.75 1.82 2.40 5.97 29.62 22.54 1.89
11.32 59.39 1.60 2.37 10.07 53.75 13.04 4.02
10.40 31.29 2.15 2.75 6.85 38.58 19.32 2.66
11.35 46.87 1.81 2.51 9.15 40.13 14.88 3.12
10.83 43.39 1.91 2.33 7.85 37.81 16.94 3.68
Maximum torque can be considered as a measure of stock modulus [118]. The torque difference (MH–ML) which shows the extent of cross linking, [119] is found to have higher variation from one compound to other. The lesser torque difference is found for the compound ‘B’ and ‘D’ which contain less gel fraction [120] (Table 2) compared to other rubber vulcanizates, which reduces the maximum rheometric torque. The obtained cross-link density values from Flory-Rehner equation [121] correspond with the variation in torque differences. The lower cross-link density in 10 wt% of ENR/nanoclay, which contains 3.09 wt% of nanoclay for ‘B’ hinders the formation of chemical cross-links and physical cross-links are formed by the clay bundles [122]. As a result of this, the decrease in MH–ML value is observed. By the use of semi reinforcing type of carbon black the scorch time reduces. This decrease in scorch time was due to presence of active cross-linking sites in the vulcanized rubber [123]. Faster cure rate index is observed in Table 2 for the compounds containing 25 wt% of SBR and 75 wt% of NR. The decrease in cure rate may be due to the greater thermal history formed during mixing, as a result of their higher compound viscosities. Also, the possible formation of a Zn complex in which sulfur and ammonium modifier participate may facilitate for the increase in rate of cure [124].
1.10.2 XRD Analysis X-ray diffraction patterns of epoxidized natural rubber with 100 wt% of nanoclay loading are shown in Fig. 5. The d-spacing (spacing between the planes in the atomic lattice) values were calculated using the Bragg’s Law. The organically modified nanoclay patterns showed an intense peak around 2h = 3.144, corresponding to the basal spacing of 2.58 nm (d001). The EC pattern showed that the d001 main diffraction peak shifted towards the lower angle 2h = 2.12, corresponding to the basal spacing of 3.79 nm (d001). The peak shift to a lower angle corresponds to the increased distance between interlayers. The higher d-spacing value is observed with ENR which signified an intercalated structure, suggesting that rubbery polymer was incorporated into the interlayer spacing. XRD data showed that the extent of intercalation was a function of the polarity of the rubber.
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Fig. 5 XRD study of ENR and nano clay
1.10.3 Mechanical Properties of the Rubber Samples Tensile strength, modulus, elongation at break, tear strength for all the compounds are shown in Table 3. The tensile strength of ‘A’, ‘E’ and ‘F’ is higher in the system, because in the case of rubber vulcanizates, the rubber chains orient themselves in the direction of stretching creating crystallites. These crystallites tie together a large number of network chains and contribute to high tensile strength and elongation. But for ‘E’ sample, the tensile strength increases with carbon black ISAF N234, possibly due to the outstanding reactivity of the carbon black acting as filler, thus enhancing the properties of the samples. While for other blend types, the tensile and tear strength starts to decrease, the filler is uniformly dispersed in the natural rubber matrix which can be attributed to the aggregation of clay nanolayers [125], also confirmed by the SEM images shown in Fig. 7. The aggregation leads to the formation of weak points in the NR matrix, accordingly reducing the elastomeric strength [126, 127]. The filler has high aspect ratio which leads to improved interfacial bonding and form filler-rubber interactions because of the high specific surface area of the filler. The mechanical properties of rubber vulcanizates markedly depended on the number of conjugate double bonds Table 3 Mechanical properties of the blends Elongation 100% Sample Tensile at break (%) modulus code strength (MPa) (MPa)
300% modulus (MPa)
Tear strength (N/mm)
Hardness (Shore A)
A B C D E F
6.61 5.01 3.45 5.66 7.77 6.32
50.6 25.2 19.7 32.4 33.8 36.6
61 57 64 69 78 70
12.19 9.44 8.55 10.19 13.17 12.65
577 457 505 526 529 524
2.25 2.01 1.85 2.44 3.01 2.89
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(sample ‘E’ and ‘F’). These observations suggest that more SBR reacts with the carbon–carbon double bonds, slower is the reversion reaction rate and hence increases the mechanical properties of the vulcanizates. The tear resistance of elastomers is mainly dependent on the processes by which stress dissipation near the tip of the growing crack takes place. Several processes such as slippage or breakage of crosslinks or chain entanglements or arresting of the growing crack by filler particles take place during the tear failure of elastomers [128]. The tear strength for all the samples is moderate except sample ‘A’, but it varies with varying the matrix ratio. So the system with carbon black SAF N110 and SRF N774 gives better reinforcing effect as well as tear strength. It is believed that at lower filler content, the filler can be dispersed well in the rubber matrix and the filler can extend further propagation. However, at higher filler content, the filler tends to form agglomerates, thus decreases the tear properties of composite. The modulus of all the NR vulcanizates increased with increasing concentration of SBR. This was because of the following possible reasons: the restriction of molecular chain mobility, and an increase in the cross-link density. The maximum torque (MH) is generally correlated with the durometer hardness and modulus. This indicates that the incorporation of organoclay filler increases the stiffness of the rubber. The hardness of ‘D,’ ‘E’ and ‘F’ rubber sample was higher than other blends as shown in Table 3. The increase in hardness of those rubber samples probably increases the cross-link density. 1.10.4 Thermal Analysis High temperature Thermal Analysis (TGA) (50–650C) curves for the sample are shown in Table 4. The temperature for the onset of degradation (T1), the temperature at which 10% degradation occurred (T10), the temperature at which 50% degradation occurred (T50) and the temperature at which 90% degradation occurred (T90) were calculated from the TGA plots. It was observed that the onset degradation temperature was more or less same for all the samples except sample ‘B’. The onset degradation temperature thereby probably decreased in the case of rubber sample due to a decrease in cross link density (Table 4). Cross linking increased the rigidity of the system, which in turn increased the thermal stability [129, 130]. This proves that increasing percentage of SBR content is responsible for the increase in thermal stability. Table 4 Thermal properties of the rubber compounds (C)
Sample code
T10
T50
T90
A B C D E F
315 307 315 317 315 318
339 339 339 410 414 413
558 558 552 549 540 558
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Fig. 6 TEM pictograph of ENR and nanoclay
1.10.5 TEM Study TEM is the most lineal method to observe the dispersion of nanoclay. TEM image of ENR with nanoclay composites is shown in Fig. 6. TEM photomicrographs showed uniform distribution of the organo-clay in the NR-SBR matrix. Nanoclay clusters were observed from the image. The dark lines are the silicate layers, in which bulks of the nanoclay dispersion are in the intercalated state. The clay structures did not break down during mixing in the NR-SBR matrix and intercalation was observed from the TEM photomicrographs. It reveals that nanoclay dispersion is in the intercalated state, which affirms the better dispersion of nanoclay in ENR [124]. It should also be noted that TEM showed that the clay layers were dispersed in the NR matrix at the nano level but XRD indicated that there were some nonexfoliated MMT layers in the NR matrix. 1.10.6 SEM Study The tensile fracture samples were scanned after gold coating, and are represented in Fig. 7. The smooth fracture surfaces and smooth filler dispersion and unidirectional tear path oriented along the direction of flow, which is smooth rubbery in nature [131], were observed for all rubber samples. Extended nanoclay platelets which are partially wrapped by the matrix due to the adsorption of the polymer on nanoclay with some tear line in branching were observed. The micrographs of the entire rubber sample are characterized by a smooth, rubbery failure (which is a smooth failure in the case of rubber samples without the formation of necking)
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Fig. 7 SEM images of different types of blends
where the additives are clearly seen and the appearance is associated with a low tensile strength. But for ‘A’, ‘E’ and ‘F’ fatigue, intermolecular and ductile type of failure was clearly observed. In ‘F’, many holes compared with the fracture surface of other samples were noticed. This hole formation may be assigned to low rubberfiller interaction as a result of detachment of the filler from the natural rubber [132]. Such holes could act as initial flaws leading to localised stress concentration during deformation. Finally, premature failure of the rubber compound occurred. This perhaps explains the reduction of both tensile and tear strength with higher filler content. Some samples like ‘E’ and ‘F’ have also shown rupture type of failure. More serious effects of hysteresis arise from chemical changes to the rubber structure at higher sustained temperatures; these effects include the rubber cross-linking, and thermal degradation leading to explosive rupture (blowout).
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This phenomenon was studied by Gent and Hindi [133]. They heated rubber specimens in a microwave oven and showed that blowout was due to the generation of gases in the interior of rubber components. 1.10.7 Du-Pont Abrader Study The mass loss of rubber against Du-Pont abrader is given in Fig. 8. The mass loss of rubber for E and F samples are lower compared to other four type blends under the same conditions. It is clear that higher abrasion resistance is mainly due to the presence of ISAF N234 type of carbon black in 65 wt% of NR, 35 wt% of SBR and 10 wt% of ENR/nanoclay in the blend system. It is also seen that samples containing 65 wt% of NR, 35 wt% of SBR and 10 wt% of ENR/nanoclay showed good abrasion property, where as SAF N110 and SRF N774 type of carbon black with 75 wt% of NR, 25 wt% of SBR and 10 wt% of ENR/nanoclay showed the high abrasion against Du-Pont abrader. 1.10.8 DIN Abrader Study Figure 9 refers to the DIN abrasion test result in terms of mass loss of rubber compounds. Compounds ‘B’ and ‘F’ showed higher abrasion resistance mainly due to the presence of ISAF N234 type of carbon black for first one and ISAF N231 type of carbon black for later. The compound C, D, and E exhibited moderate abrasion resistance property, where as SAF N110 and SRF N774 type of carbon black with 75 wt% of NR, 25 wt% of SBR and 10 wt% of ENR/nanoclay showed the high abrasion against DIN abrader. 1.10.9 Heat Buildup Study The values of heat build-up for the compounds, which showed good abrasion resistant properties against DIN and Du-Pont abrader, are shown in Table 5. For Fig. 8 Du-Pont abrasion results
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Fig. 9 DIN abrasion results
Table 5 Heat build-up of the rubber samples
Sample code
B D F
Temperature (C) Initial
10 min
20 min
50 50 50
57 59 58
62 66 64
both the compounds, the temperature development is higher, due to the presence of 40 wt% of carbon black and 10 wt% of nano clay. This may be due to the disproportionate breaking of the carbon black structure and reformation of the inter-aggregate bonds of carbon black. The compound A shows lesser heat buildup compared to B and F. The compound ‘A’ contains SAF N110 and SRF N774, whereas the compound ‘B’ and ‘F’ contain ISAF N234 high structured carbon black and ISAF N231 low structured carbon black, respectively. The use of semi reinforcing filler and 80 wt% of NR may be responsible for low heat buildup. These high temperature containing samples accelerate the fatigue of rubber components [134]. Higher tyre temperature usually means higher energy dissipation and thus higher fuel consumption [135]. Hence, it is proved that lower percentage of HSR leads to less heat build-up. It may be concluded that there is an important connection between heat build-up and the crosslinking system. Thus, the higher degree of network stability given by sulfur system generally causes high heat generation. Heat generation tests before and after aging indicate a low degree of heat build-up can be expected, even when the degree of crosslink densities is kept similar.
1.11 Summary Preparation of abrasion resistant tyre tread rubber with the help of an open tworoll-mixing mill represents a novel method for making high value rubber tyre
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tread. The NR and SBR with addition of organoclay nanocomposites obtained from this process has very good mechanical properties which can withstand the rugged working condition of automobile tyre. From this study, faster scorch time and cure time has been observed for the compounds having ISAF N231 type of carbon black. In addition, they show increase in maximum torque, which correlates with the cross-link density results. The morphology of the organoclay dispersion in ENR by solution mixing demonstrates the higher intercalation of organoclay based on the XRD results and TEM images. The FTIR study proves the interaction between ENR and organoclay. The overall mechanical properties increases for the compounds containing ISAF type carbon blacks. In DSC study, the all the samples show the same Tg except ‘E’, may be due to the ISAF type of carbon black reinforcement with 35 wt% SBR. Higher thermal stability is found for the nanocomposites containing 35 wt% SBR content. It was found that onset degradation temperature was higher for samples containing 35 wt% of SBR. From the SEM micrographs, fatigue, ductile and intermolecular fracture type of failure is clearly observed for ‘A’, ‘E’ and ‘F’, respectively. Samples containing ISAF N234 type of carbon black show higher abrasion resistance property against Du-Pont and DIN abrader. Also, sample containing 25 wt% SBR with ISAF N234 type of carbon black shows the lowest heat generation among all the samples.
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87. Grosch, K.A., Schallamach, A.: Tire friction on wet roads. Rubber Chem. Technol. 49, 862 (1976) 88. Pal, K., Pal, S.K., Das, C.K., Kim, J.K.: Influence of fillers on NR/SBR/XNBR blends. Morphology and wear. Tribol. Int. 43(8), 1542–1550 (2010) 89. Thomas, A.G.: Factors influencing the strength of rubbers. J. Polym. Sci. Polym. Symp. 48, 145 (1974) 90. Kragelskii, I.V., Nepomnyashchil, E.F. In: James, D.I. (ed.) Abrasion of Rubber. Maclern and Sons Ltd., London, Chapter 3 (1967) 91. Schallamach, A.: A theory of dynamic rubber friction. Wear 6(5), 375 (1963) 92. Nayek, S., Bhowmick, A.K., Pal, S.K., Chandra, A.K.: Wear behavior of silica filled tire tread compounds by various rock surfaces. Rubber Chem. Technol. 78, 705 (2005) 93. Kragelskii, I.V., Nepomnyashchil, E.F.: In: James, D.I. (ed.) Abrasion of Rubber. Maclern and Sons Ltd., London, Chapter 3 (1967) 94. Schnumann, R., Warlow-Davies, E.: The elastomeric component of the force of sliding friction. Proc. Phys. Soc. 54(1), 14–27 (1942) 95. Pal, K.: Speciality elastomer blends for abrasion resistant tyre tread of dump-trucks. Ph.D Thesis, IIT Kharagpur, India (2009) 96. Bhowmick, A.K.: Ridge formation during the abrasion of elastomers. Rubber Chem. Technol. 55, 1055 (1979) 97. Southern, E., Thomas, A.: Studies of rubber abrasion. Rubber Chem. Technol. 52, 1008 (1979) 98. Medalia, A.I., Alesi, A.I., Mead, J.L., Simonean, R.: Paper no. 34, Rubber Division, ACS, Cincinnati, Ohio, October, 18–21, 1988; abstract in Rubber Chem Technol 62, 165 (1989) 99. Viswanath, N., Bellow, D.G.: Development of an equation for the wear of polymers. Wear 181–183, 42 (1995) 100. Rymuza, Z.: Wear in polymer micro-pairs. Proceedings of 3rd international conference on wear of materials 125 (1981) 101. Ratner, S.B.: Connection between the wear resistance of plastics and other mechanical properties. Sov. Plast. 7, 37 (1964) 102. Lewis, R.B.: Predicting the wear of sliding plastic surfaces. Mech. Eng. 86, 32 (1964) 103. Rhee, S.K.: Wear equation for polymers sliding against metal surfaces. Wear 16, 431 (1970) 104. Lancaster, J.K.: Friction and wear. In: Jenkins, A.D. (ed.) Polymer Sciences. North-Holland Publishing Co., Amsterdam, Chapter 14 (1972) 105. Atkinson, J.R., Brown, K.J., Dowson, D.: The wear of high molecular weight polyethylene: part I: the wear of isotropic polyethylene against dry stainless steel in unidirectional motion. Trans. ASME J. Lubr. Technol. 100, 208 (1978) 106. Schallmach, A.: Abrasion pattern on rubber. Rubber Chem. Technol. 26, 230 (1953) 107. Kragelsky, I.V., Nepomnyashchil, E.F.: Friction wear of polymers. Khimiya 5 (1964) 108. Klitenik, G.S., Ratner, S.B.: Friction wear of polymers. Khimiya 77 (1964) 109. Brodsky, G.I.: Comprehensive evaluation of cord-to-rubber adhesion. Rubber World 190(5), 29–39 (1984) 110. Brodsky, G.I., Reznikovsky, M.M., Sizikov, N.N.: Rezina konstruktsionnyi material sovremennogo machinostroeniya. Khimiya 118 (1967) 111. Reznikovskii, M.M.: In: James, D.I. (ed.) Abrasion of Rubber. Maclaren, London, p. 41 (1967) 112. Rudakov, A., Kuvshinskii, E.: The mechanism of abrasion of vulcanized rubber. Rubber Chem. Technol. 37, 291 (1964) 113. Pal, K., Das, T., Rajasekar, R., Pal, S.K., Das, C.K.: Wear characteristics of styrene butadiene rubber/natural rubber blends with varying carbon blacks by DIN abrader and mining rock surfaces. J. Appl. Polym. Sci. 111, 348 (2009) 114. Dayantis, J.: The effect of pressure on the determination of the Flory-Huggins v parameter by vapour pressure measurements. Polymer 33(1), 219–221 (1992) 115. Bhatnagar, S.K., Banerjee, S.: Viscosity and molecular weight of masticated styrenebutadiene rubber. Rubber Chem. Technol. 38, 961 (1965)
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116. Ishida, H., Miller, J.D.: Substrate effects on the chemisorbed and physisorbed layers of methacryl silane-modified particulate minerals. Macromolecules 17, 1659 (1984) 117. Maiti, S.N., Mohapatro, P.K.: Mechanical properties of i-PP/CaCO3 composites. J. Appl. Polym. Sci. 42, 3101 (1991) 118. Teh PL, T.: Effects of epoxidized natural rubber as a compatibilizer in melt compounded natural rubber–organoclay nanocomposites. Eur. Polym. J. 40, 2513 (2004) 119. Sezna, J.A., Pawlowski, H.A., DeConinck, D.: New test results from rotorless curemeters. Proceeding of 136th meeting of the ACS-rubber division (1989) 120. Yehia, A.A., Ismail, M.N., Hefny, Y.A., Abdel-Bary, E.M., Mull, M.A.: Mechano-chemical reclamation of waste rubber powder and its effect on the performance of NR and SBR vulcanizates. J. Elasto Plast. 36, 109 (2004) 121. Manik, S.P., Banerjee, S.: Determination of chemical cross-links in rubbers. Die Angewandte Makromolekuiare Chemie 6, 171 (1979) 122. Zhu, L., Wool, R.P.: Nanoclay reinforced bio-based elastomers: synthesis and characterization. Polymer 47, 8106 (2006) 123. Rajasekar, R., Pal, K., Heinrich, G., Das, A., Das, C.K.: Development of NBR-nanoclay composites with epoxidized natural rubber as compatibilizer. Mater. Des. 30, 3839 (2009) 124. De, D., Maiti, S., Adhikary, B.: Reclaiming of rubber by a renewable resource material (RRM). III. Evaluation of properties of NR reclaim. J. Appl. Polym. Sci. 75, 1493 (2000) 125. Agag, T., Koga, T., Takeichi, T.: Studies on thermal and mechanical properties of polyimide–clay nanocomposites. Polymer 42, 3399 (2001) 126. Nielsen, L.E. (ed.): Mechanical Properties of Polymers and Composites. Marcel Dekker, New York, Chapter 2 (1974) 127. Sharifa, J., Yunus, W.M.Z.W., Dahlan, K.Z.H.M., Ahmad, M.H.: Preparation and properties of radiation crosslinked natural rubber/clay nanocomposites. Polym Test 24(2), 211–217 (2005) 128. Thomas, S., Kuriakose, B., Gupta, B.R., De, S.K.: Scanning electron microscopy studies on tensile, tear and abrasion failure of plasticized poly (vinyl chloride) and copolyester thermoplastic elastomers. J. Mater. Sci. 21, 711 (1986) 129. Maity, M., Khatua, B.B., Das, C.K.: Effect of processing on the thermal stability of the blends based on polyurethane: part IV. Polym. Degrad. Stab. 72, 499 (2000) 130. Gann, R.G., Dipert, R.A., Drews, M.J.: Flammability. In: Kroschwitz, J.I. (ed) Encyclopedia of Polymer Science and Engineering, 2nd edn. John Wiley & Sons, Inc., New York, 7:154– 210 (1985) 131. Pal, K., Das, T., Pal, S.K., Das, C.K.: Use of carboxylated nitrile rubber and natural rubber blends as retreading compound for OTR tires. Polym. Eng. Sci. 48, 2410 (2008) 132. Siriwardena, S., Ismail, H., Ishiaku, U.S.: A comparison of white rice husk ash and silica as fillers in ethylene–propylene–diene terpolymer vulcanizates. Polym. Int. 50, 707 (2001) 133. Gent, A.N., Hindi, M.: Heat build-up and blowout of rubber blocks. Rubber Chem. Technol. 63, 892 (1988) 134. Medalia, A.I.: Heat generation in elastomer compounds: causes and effects. Rubber Chem. Technol. 64, 481–492 (1991) 135. Park, D.M., Hong, W.H., Kim, S.G., Kim, H.J.: Heat generation of filled rubber vulcanizates and its relationship with vulcanizate network structures. Eur. Polym. J. 36, 2429–2436 (2000)
Elastomer Clay Nanocomposites for Packaging V. Mittal
Abstract Packaging application of the materials require strong barrier as well as mechanical properties. Majority of the reports in the literature have studied the mechanical properties while neglecting the barrier properties and it is generally assumed that the improvement in mechanical properties leads to automatic enhancement of barrier performance. However, it may not be true in all the cases as the permeation properties are more sensitive to interface between the polymer and filler phases. The incompatibility between the filler surface modification and polymer or the presence of excess surface modification on the filler surface can negatively impact the permeation properties, while the mechanical performance may still enhance. The current chapter focuses on the two very important elastomeric materials used in large amounts in packaging industry. Various factors affecting the barrier performance have been quantified. Use of not only montmorillonite, but also vermiculite has been presented. The importance of clean filler surface on the composite properties has also been underlined. The determination of average aspect ratio as a tool to control or tune the microstructure and properties of the nanocomposites has been shown with the help of finite element computer models.
V. Mittal (&) Institute of Chemical and Bioengineering, Department of Chemistry and Bio-Engineering, ETHZ Zurich, 8093, Zurich, Switzerland e-mail:
[email protected] Present Address: V. Mittal Polymer Research, BASF SE, 67056, Ludwigshafen, Germany V. Mittal Chemical Engineering Program, The Petroleum Institute, 2533, Abu Dhabi, United Arab Emirates
V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_9, Ó Springer-Verlag Berlin Heidelberg 2011
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Keywords Gas barrier Mechanical properties Delamination Clay Aspect ratio Intercalation Degradation Interface Polyurethane Epoxy Montmorillonite Vermiculite Finite element models
1 Introduction Elastomeric nanocomposites form a significant proportion of all the nanocomposites synthesized in general. As is common with polymer nanocomposites based on polypropylene, polyethylene, etc., majority of the studies on elastomeric nanocomposites also report enhancements in the mechanical performance of the polymers. The other properties of immense importance like gas barrier properties have been generally neglected. It is commonly believed that the improvements in the mechanical properties lead to automatic reinforcement in other properties too, but it may not be true in all the cases. The reason for such difference is the significant dependence of the gas barrier properties on the interface development between the organic and inorganic phases at nanoscale. Any minor incompatibility at the interface may not affect the mechanical properties, but it significantly affects the gas barrier properties as the incompatibility at interface may lead to a source of increased diffusion though the materials thus negating the whole purpose of polymer reinforcement with the high aspect ratio inorganic clay platelets. Polymers find increasing use in packaging applications where the polymer packaging materials are required to have high strength and high barrier to gas (oxygen, water vapor, carbon dioxide, aroma, etc.) among other properties. Thus, it is important to study the gas barrier properties of the elastomeric composites to establish the factors which affect the barrier performance of these nanocomposites along with the mechanical performance enhancement. These factors vary with the nature of the polymer, nature of the surface modification of the inorganic filler as well as processing conditions used for the synthesis of the nanocomposites. Understanding these factors can lead to optimal design of the high barrier (and high strength) organic–inorganic hybrid materials for packaging applications. The chapter would focus on elastomer polyurethane and epoxy polymers for the synthesis of nanocomposites to achieve high gas barrier materials owing to their extensive use in the packaging industry mostly as adhesives. Figure 1 shows an simplified representation of a packaging laminate where the poly(ethylene terepthalate) and polypropylene foils are adhered together by the use of an adhesive comprising of either epoxy or polyurethane [1]. The foils are chosen by the virtue of Fig. 1 Representation of a commercial packaging laminate. Reproduced from [1] with permission from Nova Science Publishers
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the properties expected from them, but the choice of the epoxy or polyurethane based adhesive is solely to bond together the other foils and it does not contribute towards the barrier performance of the laminate. By the inclusion of the layered silicate fillers with high aspect ratio crystalline platelets, the adhesive can also be functionalized to contribute towards the barrier performance of the laminate and by virtue of thus functionality, the thickness of the laminate can be reduced or one or two foils of the laminate can be avoided thus leading to significant material savings. The elastomeric nanocomposites thus can also be developed as free standing barrier foils for packaging.
2 Polyurethane Clay Nanocomposites Polyurethane rubbers or elastomers have many applications in foams, coatings, adhesives of packaging laminates, fibers, etc. Use of polyurethane for the packaging applications requires the enhancement of both barrier properties as well as mechanical performance for which the plate like clay particles are added to the matrix as the thin platelets lead to efficient stress transfer between the organic and inorganic phases, increase interfacial contacts between the organic and inorganic phases as well as increase the mean path length of the permeant molecules. Thus the permeant molecules have to wiggle around the platelets in a random walk caused by the tortuouity in the path of the permeant molecules owing to the incorporation of clay platelets. In one such representative study, elastomeric polyurethane nanocomposites were studied by incorporating montmorillonite modified with different surface modifications [2]. The chemical modifications exchanged on the surface of the montmorillonite were bis(2-hydroxyethyl) hydrogenated tallow ammonium (Nanofil 804), alkylbenzyldimethylammonium (benzalkonium, Nanofil 32), and dimethyl dihydrogenated tallow ammonium (Nanofil 15). The different surface modifications were chosen in order to study the effect of the chemical structure of the modification and subsequent interactions with the polymer on the composite microstructure and properties. Especially the modification with hydroxyethyl moieties was expected to have much better polarity match with the polymer than the other modifications. The nanocomposites were generated by solvent intercalation approach and nanocomposite coatings were drawn on suitable substrates. Figure 2a shows the X-ray diffractograms of the Nanofil 804 filled polyurethane nanocomposites with different filler volume fractions. The diffraction pattern of the pure clay has also been shown. As is clear from the diffractograms, the diffraction peak of the pure filler was shifted to lower angles in the composites thus indicating intercalation of the polymer in the filler interlayers. However, the presence of the diffraction peak in all the composites indicates that the filler was not completely delaminated. Also, the same peak position was observed in the composites irrespective of the filler amount indicating the interlayers were intercalated by similar amount of polymer. As the intensity of the X-ray diffraction
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Fig. 2 a X-ray diffractograms of the polyurethane nanocomposites containing different volume fractions of the Nanofil 804 modified montmorillonite and b TEM micrograph of the polyurethane nanocomposites comprising of 2.86 vol% of the Nanofil 804 modified montmorillonite. Reproduced from [2] with permission from American Chemical Society
peaks is qualitative in nature owing to the dependence on sample preparation, filler impurities, orientation, etc., the composite microstructure was also examined with transmission electron microscopy as shown in Fig. 2b. The clay platelets in the composites with 2.86 vol% of the Nanofil 804 filler were observed to be extensively exfoliated. Other volume fractions of the Nanofil 804 filler in the composites led to the generation of similar morphology. The composites generated with other fillers like Nanofil 32 and Nanofil 15 on the other hand had more intercalated morphology than exfoliated. The increase in the basal plane spacing observed in the X-ray diffractograms was also the maximum in the case of Nanofil 804 composites, even though the other fillers had initial higher basal spacing. The interaction of the hydroxyl groups on the surface of the filler platelets were expected to have caused better interactions with the polyurethane polymer thus leading to higher extents of filler exfoliation. Even a chemical reaction between the two can be expected thus chemically tethering the polymer chains with the filler platelets, but it could not be confirmed spectroscopically.
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Fig. 3 Relative a oxygen and b water vapor permeation through the polyurethane nanocomposites as a function of filler volume fraction for different fillers. Reproduced from [2] with permission from American Chemical Society
Usability of the generated nanocomposites for the packaging applications was tested by measuring their barrier performance. Figure 3 shows the oxygen and water vapor permeation through the nanocomposites as a function of filler volume fraction. The oxygen permeation in the composites with fillers of Nanofil 804 and Nanofil 32 decreased as a function of filler volume fraction. 30% reduction could be achieved at a filler fraction of 3 vol%. Though the Nanofil 804 filler was expected to chemically bind with the polymer thus improving the barrier performance, however, the oxygen permeation behavior of Nanofil 804 filled composites was similar to Nanofil 32 filled composites. It indicated that either the Nanofil 32
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filler owing to the presence of benzyl rings in the surface modification also had a better interaction with the polymer or the chemical reaction in the case of Nanofil 804 filler with the polymer did not take place. More surprising was the result obtained for Nanofil 15 filled composites, where the permeation was observed to increase with increasing the filler volume fraction. As the only difference in the three different systems lies in the nature of the surface modification ionically bound to the filler platelets, therefore, it is clear that the chemical architecture of the surface modification did affect the composite microstructure and properties. The modification in the case of Nanofil 15 is completely non-polar which may not mix well with the polar polymer chains thus causing incompatibilities at the interface. These incompatibilities are not detected in X-ray or microscopy but were observed to affect the interface sensitive barrier properties. The incompatibility can lead to the generation of voids or areas of increased free volume which are responsible for the increase in the permeation. It also explains the increase of permeation as a function of filler volume fraction. The water vapor permeation behavior through the composites is also depicted in Fig. 3b. Here the permeation decreased in all the three cases as a function of filler volume fraction and there was no increase in the case of Nanofil 15. It is owing to the different mode of transport of water molecules through the polymer matrix. Water molecules interact with each other as well as with the polymer matrix thus forming big clusters which cannot pass through the voids generated by the mismatch between the polymer and filler surface modification. Thus, the water vapor permeation is more of a function of hydrophobicity of the filler. More hydrophobic the filler is, higher the decrease in the water vapor permeation through the composites. Similarly, other studies have highlighted the use of polyurethane composites for packaging applications. Relative water vapor permeation of a poly(urethane urea)dimethyl dehydrogenated tallow ammonium modified montmorillonite composite was reported to show a decrease of five times at a 6 vol% filler concentration [3]. Oxygen permeation through polyurethane nanocomposites was observed to decrease by 50 or 15% at 4 wt% of montmorillonite surface treated with hexadecylammonium or dodecyltrimethylammonium salts, respectively, and by 35% at 4 wt% Cloisite 25A (dimethyl dehydrogenated tallow, 2-ethylhexylammonium) [4]. Similarly, water vapor permeation of polyurethane was reported to have significant decrease up to 20 wt% organically modified montmorillonite concentration by Tortora et al. [5]. Mechanical performance of the polyurethane nanocomposites has been reported in many studies [6–13]. Figure 4 shows the improvement in the tensile modulus of the polyurethane nanocomposites as a function of filler fraction [6]. The modulus increased more than twice the value of pure polymer at a filler loading of 10 wt%. Apart from tensile modulus, both tensile strength as well as strain at break were also reported to significantly improve as a function of filler fraction. The authors suggested that the improvements in tensile modulus and strength can be attributed to the reinforcing effect of dispersed clay platelets whereas improvement in the strain at break can be a result of the plasticizing effect of the ammonium ions ionically bound in the filler interlayers.
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Fig. 4 Tensile modulus of the polyurethane nanocomposites as a function of filler weight fraction. Reproduced from [6] with permission from American Chemical Society
Fig. 5 Force-extension curves for the pure polymer as well as polymer nanocomposites containing different vol% of the filler. Reproduced from [3] with permission from American Chemical Society
Xu et al. [3] also reported force-extension curves for the poly(urethane urea) based nanocomposites as shown in Fig. 5. The authors reported a significant increase in modulus and strength of the composites owing to the nano dispersed filler platelets. It was observed that for 20 wt% composite, the modulus and tensile strength increased by more than 300 and 30%, respectively. Apart from that, there was no negative impact on the ductility of the composites as the elongation to break of the 20 wt% nanocomposite was observed to increase by 50% as compared to pure polymer.
3 Epoxy-Montmorillonite Nanocomposites Similar to the polyurethane nanocomposites, elastomeric epoxy nanocomposites were also synthesized in order to present their use for packaging applications.
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Fig. 6 Chemical structures of the filler modifications as well as epoxy–amine system used fro the synthesis of epoxy nanocomposites. Reproduced from [14] with permission from American Chemical Society
The interaction of filler surface modification with the epoxy polymer was also similarly carried out by using various surface modifications as shown in Fig. 6 [14]. The modifications used were benzyldibutyl(2-hydroxyethyl)ammonium chloride (Bz1OH), benzylbis(2-hydroxyethyl)butylammonium chloride (Bz2OH), benzyltriethanolammonium chloride (Bz3OH), and benzyl(2-hydroxyethyl)methyloctadecylammonium chloride (BzC18OH), benzyldimethylhexadecylammonium chloride (BzC16) and dioctadecyldimethylammonium chloride (2C18). The modifications differed in extent of polarity by the virtue of different chemical structures. The modifications with one, two and three hydroxyl groups per molecule were exchanged on the surface in order to observe the effect of the hydroxyl groups on the resulting composite microstructure as well as properties. Similarly, modifications also included benzyl groups, combination of benzyl groups with hydroxyl groups and octadecyl chains. Apart from that, modification with only octadecyl chains was also used. Similar to the case of polyurethane nanocomposites, the hydroxyl group containing modifications were expected to tether the polymer chains with the filler surface. Apart from that, benzyl groups in the modification molecules were expected to bring stronger van der Waals forces of interaction between the filler and the polymer. The long alkyl chains were used to facilitate better interlayer spacing, which may be of help in intercalating more and more polymer thus leading to the filler delamination. X-ray diffraction studies on the filler dispersions in solvent and composites revealed interesting insights into the systems. The prepolymer was observed to
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have negative interaction with I, II and III surface modifications owing to mismatch of polarity. The dispersions of the filler in solvent were observed to have lowering of the basal plane spacing after the addition of the prepolymer to the dispersion. The diffractogram in the case of benzyldimethylhexadecylammonium modified clay suspension in solvent had no diffraction peak indicating complete exfoliation, but a diffraction peak at 3.55 nm appeared after the addition of prepolymer. Similarly for dioctadecyldimethylammonium modified clay, a diffraction peak for the filler suspension in solvent was observed at 3.69 nm which became 3.54 after the addition of prepolymer. For benzyl(2-hydroxyethyl)methyloctadecylammonium modified clay, a decrease from 4.12 to 3.89 nm was similarly observed. On the other hand, the modifications with hydroxyl groups (especially one hydroxyl group per molecule) had positive interaction with the prepolymer. The basal plane spacing of especially Bz1OH and Bz2OH modified fillers was not affected or increased by the addition of the epoxy prepolymer. These fillers retained the basal plane spacing values around 1.9 nm before and after the addition of prepolymer. Composite basal plane spacing values were observed to decrease in all the cases as compared to the filler suspensions in solvent owing to the evaporation of solvent during crosslinking. As the X-ray diffraction provided only qualitative insight into the system, it was further synergized with the microscopy investigations. The TEM investigation of the 3.5 vol% composites with IV and III modified montmorillonites is shown in Fig. 7. The composite with benzyldibutyl(2-hydroxyethyl)ammonium modified clay was observed to have extensive filler exfoliation, whereas, the composite with benzyldimethylhexadecylammonium modified clay was more intercalated in nature. Composites with V as filler modification were also exfoliated in nature, whereas with VI modified clay as filler, the composites had intercalated morphology probably owing to high extent of polarity of the interlayer which may also not match well
Fig. 7 TEM micrographs of the 3.5 vol% filler epoxy nanocomposites using modifications a benzyldibutyl(2-hydroxyethyl)ammonium and b benzyldimethylhexadecylammonium. Reproduced from [14] with permission from American Chemical Society
242 Table 1 Oxygen and water vapor permeation through the 3.5 vol% epoxy nanocomposites
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Permeability coefficient (oxygen) [cm3 lm/ (m2 d mmHg)]a
Water vapor transmission rate [g lm/(m2 d mmHg)]a
Neat epoxy Cloisite Na+ BzlOH Bz20H Bz30H BzC180H BzC16 2C18
2.0 1.6 0.77 0.78 1.0 2.2 1.6 3.7
10 23 6.7 5.8 7.1 5.7 5.3 6.8
Reproduced from [14] with permission from American Chemical Society a Relative probable error 5%
with the epoxy polymer. The other modifications also led to the generation of the intercalated nanocomposites. The basal plane spacing values of the fillers modified with Bz1OH as well as Bz2OH were much smaller than the modifications containing long octadecyl chains. Even then, the higher extent of filler exfoliation was observed in composites containing Bz1OH as well as Bz3OH modified fillers. It again confirms the importance of the interactions between the organic and inorganic phases in developing the composite microstructure. The initial higher basal plane spacing is not as important as matching the compatibility between the phases. The observations from X-ray diffraction as well as microscopy were also reflected in the composite properties. Table 1 shows the oxygen and water vapor barrier properties of the epoxy nanocomposites at 3.5 vol% filler fraction. A plot between the relative oxygen permeation and the basal spacing of the filler in the composite led to an indication that the permeation actually increased with the increase in basal plane spacing, which cannot be true. Thus, there is an absence of correlation between the two factors indicating that the increase in the basal plane spacing by the intercalation of polymer in the interlayers is not responsible for the improvement in the oxygen barrier. It is rather the exfoliated platelets which are more responsible towards this barrier as the aspect ratio of the platelets is fully recovered, whereas the intercalated platelets have no increase in the aspect ratio. The oxygen permeation through the composites containing BzC16 modified filler was observed to be similar as the composites with untreated montmorillonite. The composites with 2C18 as well as BzC18OH modified montmorillonites were observed to have increased oxygen permeation as compared to pure epoxy matrix. This behavior of incompatibility of the alkyl chains with polar polymer matrix confirms the X-ray and TEM findings. Similar response was also observed for the polyurethane nanocomposites. Even the presence of alkyl chain and hydroxyl groups together in the surface modification had negative influence on the microstructure and properties of the composites. Maximum decrease in the oxygen permeation was observed for the Bz1OH and Bz2OH modified montmorillonites also indicated earlier by the virtue of positive interaction of the filler modification
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with the epoxy prepolymer. Increasing the number of OH groups in the surface modification to three was not very effective as it seemed to have become too polar for the polymer matrix to mix well. However, the performance of the composites with Bz3OH modified montmorillonite was much better than the composites with montmorillonites modified with long alkyl chains as 50% reduction in the oxygen permeation as compared to pure polymer could still be achieved. The water vapor permeation through the composites is also represented in Table 1. The various surface modifications though had significant effect on the oxygen permeation properties of the nanocomposites, the water vapor permeation was not significantly affected by the different surface modifications and the decrease in the water vapor permeation was a function of the hydrophobicity of the filler also observed in the polyurethane composites. The composite with untreated montmorillonite had a massive increase in the water vapor permeation owing to its high hydrophilicity. The composites with surface modified montmorillonites had reduced hydrophilicity, thus, were also observed to have lower water vapor permeation. The composites where the oxygen permeation was observed to deteriorate were also observed to have improved water vapor resistance owing to the different modes of diffusion of water vapor through the polymer matrices than oxygen as mentioned in the case of polyurethane nanocomposites. These results thus direct the optimal design of epoxy nanocomposites for their use in packaging applications. Figure 8 also shows the oxygen and water vapor permeation through the epoxy nanocomposites as a function of filler volume fraction. Composites with benzyldibutyl(2-hydroxyethyl)ammonium or Bz1OH and benzyldimethylhexadecylammonium or BzC16 modified montmorillonites have been compared. Both the composite series were observed to have reduced oxygen permeation as a function of filler volume fraction, but the composites containing Bz1OH had much more impressive decrease than the BzC16 composites. The oxygen permeation was reduced to one fourth at 5 vol% loading of the Bz1OH modified filler, whereas in the case of BzC16 composites, the oxygen permeation decreased by roughly 20% at 5 vol% filler loading. The permeation results were also fitted with theoretical models to generate an idea about the average aspect ratio of the platelets in the composite. It was observed that the platelets in the case of BzC16 composites had an average aspect ratio of 50, whereas the platelets in Bz1OH nanocomposites had an average aspect ratio of 250–300 indicating much higher exfoliation of the filler which subsequently affects the permeation properties and hence composite applications. As the interaction of permeant with the polymer matrix defines the diffusion behavior, the water vapor permeation was different from the oxygen permeation trends as shown in Fig. 8b. The permeation through the composites containing BzC16 modified montmorillonite decrease by 50% at 5 vol% loading and the performance was also superior to the Bz1OH modified fillers. Epoxy nanocomposites with enhanced mechanical properties have been extensively reported [15–23]. In one such study, Lan and Pinnavaia [15] reported the epoxy nanocomposites with montmorillonites modified with different chain length surface modifiers. Tensile properties of the epoxy nanocomposites filled with varying amounts of octadecyl ammonium modified montmorillonite have been
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Fig. 8 a Oxygen and b water vapor barrier properties of epoxy nanocomposites as a function of filler volume fraction and filler modifications of benzyldibutyl(2hydroxyethyl)ammonium (Bz1OH) and benzyldimethylhexadecylammonium (BzC16). The dotted lines act as guide. Reproduced from [14] with permission from American Chemical Society
detailed in Fig. 9. It was observed that both tensile strength as well as tensile modulus increased nearly linearly with increasing filler content in the composite. More than ten times increase in both strength and modulus was observed at 15 wt% of the filler. Majority of the reported studies on epoxy nanocomposites have used nonreactive surface modifications or sometimes modifications with van der Waals interactions with the polymer matrix have also been employed. The modifications containing hydroxyl groups have also been exchanged on the filler surface in the hope of chemical reaction with the epoxy polymer [24], but the reaction could not be confirmed. In another interesting study, Ma et al. ionically exchanged the hardener M-xylylenediamine containing two amine groups. One reactive group of the hardener was converted into a cation and was subsequently ion exchanged on the surface of the montmorillonite platelets to generate hardener grafted clay as shown in Fig. 10. The hardener modified clay was then mixed with epoxy prepolymer and blended well followed by heating. Following this procedure, the
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Fig. 9 Tensile strength and tensile modulus of the epoxy nanocomposites filled with varying amounts of octadecyl ammonium modified montmorillonite. Reproduced from [15] with permission from American Chemical society
polymer was cured by using 4-aminophenyl sulfone (DDS) to generate completely disordered nanocomposites. The authors also compared the nanocomposites generated by using conventional modification of the filler (containing OH groups). Xray diffraction peaks were observed in the case of these composite, whereas the diffractogram had no diffraction signal in the composites generated by hardener grafting method.
4 Epoxy Vermiculite Nanocomposites Majority of the studies on nanocomposites have focused on the use of montmorillonite as the fillers. Owing to lower charge density of the montmorillonites, a larger area per cation is available on the surface of the platelets. As the crosssection area of the exchanged surface modification molecules is smaller than the
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Fig. 10 Representation of the modification of the clay with hardener and subsequent epoxy crosslinking. Reproduced from [24] with permission from American Chemical Society
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area available per cation on the surface of the montmorillonite platelets, the surface modifications do not lie straight on the surface. This leads to lower basal plane spacing in the interlayers. However, if the medium charge density minerals like vermiculite (0.5–0.8 equiv mol-1) are chosen, they can provide the potential of obtaining higher basal plane spacing for the same ammonium modification as compared to montmorillonites, the higher charge density leads to lower area per cation on the surface, thus, more amount of organic modification can be exchanged on the same area leading to more upright positioning of the molecules. However, the higher charge density minerals do not swell easily in water and complete cation exchange is generally not possible. But, an optimal exchange can still be generated if the cation exchange process is properly controlled and on occasions also repeated to enhance the extent of surface coverage with the organic modification. It was mentioned in the earlier section that it is more the interaction between the polymer and surface modification than the initial higher basal plane spacing, which is responsible for the filler exfoliation. In the case of epoxy montmorillonite nanocomposites, Bz1OH modification though generated little basal plane spacing increment in the filler, but the composites were extensively exfoliated than the case of modifications which generated initial higher basal plane spacing because of incompatibility of these modifications with the prepolymer. However, the interactions are also dependant on the nature of inorganic substrate, thus, it is of importance to study the epoxy nanocomposites with vermiculite system too along with comparison with the montmorillonite system [25]. Also, the aspect ratio of the platelets is higher in case of vermiculite than the montmorillonite which is further helpful for the application of composite materials for packaging. The vermiculite platelets in the reported study [25] were modified with surface modifications: benzyldibutyl(2-hydroxyethyl)ammonium (Bz1OH), benzyldimethylhexadecylammonium (BzC16) and benzyl(2-hydroxyethyl)methyloctadecylammonium (BzC18OH). The exchange with Bz1OH cations was however not successful and the cations did not exchange the sodium cations present on the surface of the vermiculite platelets. The other two cations BzC16 and BzC18OH could be successfully exchanged and the extent of cation exchange was also near complete. Figure 11 shows the thermogravimetric analysis of the modified vermiculite samples. The degradation of the organic matter in the temperature range of 200–350°C as one peak indicates the absence of any excess modification molecules which would degrade at lower temperatures. The presence of excess surface modification molecules in the filler interlayers can have negative impact on the composite properties, as is explained in Sect. 5. Table 2 details the X-ray diffraction analysis of the pure filler, filler suspension in solvent as well as epoxy nanocomposites with 3.5 vol% of filler fraction. Basal plane spacing of 3.40 and 3.25 nm were observed for BzC18OH and BzC16 modified vermiculite as compared to 1.22 nm for pristine vermiculite. These values were much higher than the montmorillonites similarly modified (2.06 nm for BzC18OH and 1.87 nm for BzC16 modification). The suspension the filler in the solvent led to further enhancement of the basal plane spacing, but complete exfoliation of the filler was not observed. The addition of the prepolymer to the
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Fig. 11 Thermogravimetric analysis of benzyldimethylhexadecylammonium (BzC16) and benzyl(2hydroxyethyl)methyloctadecylammonium (BzC18OH) modified vermiculite. Reproduced from [25] with permission from Sage Publishers
Table 2 Basal plane spacing of 3.5 vol% filler fraction Filler d-Spacing filler powder (nm)
the filler, solvent suspensions as well as nanocomposites with d-Spacing of filler suspended in DMF (nm)
d-Spacing of filler suspended in DMF ? epoxy (nm)
d-Spacing in epoxy composite (nm)
Na-vermiculite BzC180H BzC16
1.42 3.80 3.34
1.42 3.80 3.53
1.29 3.96 3.68
1.22 3.40 3.25
Reproduced from [25] with permission from Sage Publishers
suspensions increased the basal plane spacing further though even by the addition of prepolymer, complete exfoliation of the filler was not achieved. However, it still indicates the intercalation of prepolymer in the filler interlayers and this behavior is opposite from that observed in the case of montmorillonite system, indicating the influence of the inorganic substrate on the interactions between the components. The synthesis of the composite after curing of the polymer and removal of solvent led to further increase in the basal plane spacing indicating more polymer intercalated in the interlayers during composite synthesis. The basal plane spacing of the filler in composites filled with montmorillonites was observed to decrease after the solvent evaporation. However, the presence of the diffraction peak in the composites indicated again the absence of complete filer exfoliation. The microscopy analysis on the composites as reported in Fig. 12 for BzC16-vermiculite epoxy nanocomposites revealed the presence of mixed morphology where a part of filler was exfoliated and a part was intercalated to varying extents. Barrier performance of the nanocomposites revealed interesting differences in comparison with the montmorillonite nanocomposites (Table 3). The oxygen permeation of the epoxy matrix is quite low indicating further decrease in this value is not very straightforward. The composites with BzC18OH and BzC16 filler modifications were observed to have oxygen permeation values of 1.5 and 1.4 cm3/
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Fig. 12 TEM micrograph of the 3.5 vol% BzC16-vermicukte epoxy nanocomposite. Reproduced from [25] with permission from Sage Publishers
Table 3 Oxygen and water vapor permeation through the 3.5 vol% vermiculite epoxy nanocomposites Composite Oxygen permeability Water vapor coefficienta [cm3 lm/ transmission ratea [g lm/(m2 d mmHg)] (m2 d mmHg)] Neat epoxy Na-vermiculite BzC18OH BzC16
2.0 1.7 1.5 1.4
10.0 37.0 7.5 9.7
Reproduced from [25] with permission from Sage Publishers a Relative probable error 5%
m2 d mmHg respectively indicating further decrease in the oxygen permeation of epoxy could be obtained by the nanoscale reinforcement of polymer with vermiculite platelets. The composites with same modifications for the montmorillonite system had oxygen permeation values of 2.2 and 1.6 cm3/m2 d mmHg respectively indicating that the vermiculite system had better polymer filler interactions than the montmorillonite system. The water vapor permeation of the composites was less affected owing to the residual polarity of the vermiculite fillers after cation exchange.
5 Effect of Excess Surface Modification The surface modification of the filler platelets is a very important factor which significantly affects the generation of composite microstructure and hence the
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Fig. 13 Representation of presence of excess surface modification molecules on the filler surface. Reproduced from [26] with permission from Springer
resulting composite properties. During the surface modification, generally the modification in excess of the cation exchange capacity of the mineral is added to ensure complete cation exchange. However, the removal or washing of this excess of surface modification molecules is not very simple and a number of washing cycles are necessary to obtain completely clean filler [26]. Commercially treated organo-montmorillonites have been generally observed to contain excess of modifier molecules. Figure 13 shows the representation of the presence of excess surface modification molecules on the filler surface [26]. The excess molecules form a pseudo bilayer and get trapped between the ionically bound molecules. A number of studies have reported that the presence of excess and unattached molecules of surface modification on the filler surface can have a detrimental effect on the microstructure and properties of the composites [27–29]. Two different phenomena by which these excess molecules may affect the composite microstructure and properties have been observed: firstly, the unattached molecules present as pseudo bilayer degrade thermally at lower temperature and the resulting products can further degrade the polymer matrix thus resulting in the deterioration of the interface between the polymer and filler. This point is more of concern when high temperature compounding of the filler with the polymer is involved. Secondly, the excess modification molecules can also specifically interact with the polymer even at lower or room temperature. In case of epoxy system, the free ammonium head groups can bond with the epoxy prepolymer thus inducing system instability and disturbing filler polymer interfacial interactions. Such interactions specifically impact the interface sensitive permeation properties, thus, it is important to ensure the clean filler surface in order to eliminate any deteriorating impact on these properties. To study the effect of excess surface modification molecules on especially the barrier properties, elastomer epoxy nanocomposites with dioctadecyldimethylammonium (2C18), bis(2-hydroxyethyl) methylhydrogenatedtallowammonium (C182OH) and benzylhexadecyldimethylammonium (BzC16) modified montmorillonites were synthesized. Montmorillonites with two different cation exchange capacities of 680 and 880 leq g-1 were used. Figure 14 shows the thermogravimetric analysis of commercially procured and self modified BzC16M680 and 2C18M880 montmorillonites. The presence of two thermal degradation peaks
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Fig. 14 Thermogravimetric analysis of a commercially modified and b self modified BzC16M680; c commercially modified and d self modified 2C18M880. Reproduced from [26] with permission from Springer
between 200 and 300°C in Fig. 14a, c indicates the presence of excess surface modification molecules whose degradation occurs at lower temperature than the molecules which are ionically bound to the filler surface. It was further observed that roughly 15–20% of the total weight loss can be attributed to this excess material indicating a significant amount of excess in the commercially modified montmorillonites. The self modified montmorillonites were washed rigorously following multiple cleaning cycles and the thermogravimetric analysis of two of these montmorillonites is depicted in Fig. 14b, d. The absence of any low temperature thermal degradation peak or the presence of only one degradation peak in the region of 200–300°C confirmed the absence of any excess surface modification molecules indicating the excess of surface modification molecules used during the exchange reaction could be successfully washed off. Table 4 shows the oxygen permeation through the 3.5 filler vol% epoxy nanocomposites synthesized by using commercially modified clays. A value of 2 cm3 lm/(m2 d mmHg) was observed for pure epoxy which was observed to increase to high values in all the composites indicating the negative impact of
252 Table 4 Oxygen permeation through the 3.5 vol% epoxy nanocomposites synthesized with commercially modified montmorillonites
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Oxygen permeability coefficient [cm3 lm/(m2 d mmHg)]
Neat epoxy 2C18M680 C182OHM680 BzC16M680 BzC16M880 2C18M880 C182OHM880
2.0 4.8 5.9 2.9 2.8 4.6 5.5
Reproduced from [26] with permission from Springer
Table 5 Oxygen permeation through the 3.5 vol% epoxy nanocomposites synthesized with self modified or washed montmorillonites
Composites
Oxygen permeability coefficient [cm3 lm/(m2 d mmHg)]
Neat epoxy BzC16M680 2C18M680 BzC16M880 2C18M880 C182OHM880 (washed)
2.0 1.8 3.9 1.6 3.7 1.5
Reproduced from [26] with permission from Springer
excess surface modification on the composite properties. Also, no effect of the chemical architecture of the surface modification molecules as well as cation exchange capacity of the montmorillonite on the oxygen permeation properties was observed. Especially high was the oxygen permeation through the C182OHM680 and C182OHM880 filled composites where the oxygen permeation values were enhanced to 5.9 and 5.5 cm3 lm/(m2 d mmHg) respectively. Table 5 describes the oxygen permeation of the composites when self treated or washed montmorillonites were used as filler. Completely different oxygen permeation behavior through the composites was observed as compared to commercially modified systems. Oxygen permeation of 1.8 cm3 lm/(m2 d mmHg) instead of 2.9 cm3 lm/(m2 d mmHg) was observed for BzC16M680 system. Similarly, oxygen permeation of 1.6 cm3 lm/(m2 d mmHg) was observed for BzC16M880 instead of 2.8 cm3 lm/(m2 d mmHg) for commercially modified system. Composites with 2C18 modified fillers though also saw a decrease in the oxygen permeation when self modified montmorillonites were used, but it was expected as already mentioned in the earlier section owing to mismatch between the polymer and filler modification. Remarkable improvement in the oxygen permeation behavior of C182OHM880 system was observed after washing. Before washing, very high oxygen permeation of 5.5 cm3 lm/(m2 d mmHg) was observed, which was reduced to 1.5 cm3 lm/(m2 d mmHg) after washing the montmorillonites. These results clearly establish the importance of filler cleanliness on the resulting composite properties. The higher cation exchange capacity of
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the montmorillonite was also observed to perform better than the low capacity mineral. The more polar surface modification was also observed to have better oxygen resistance properties owing to better interaction with polymer and hence higher extent of filler exfoliation. Another important point to mention here is that though the presence of excess of surface modification molecules in the filler interlayers impacted the barrier properties significantly, but the composites with commercially modified montmorillonites or self treated montmorillonites were observed to have no difference in the microstructure as evaluated with X-ray and microscopy thus further underlining the sensitive nature of permeation properties.
6 Aspect Ratio and Gas Permeation As mentioned above, the increase in the aspect ratio of the filler in composite represents its exfoliation and it is only the exfoliated (and not intercalated) platelets which have maximum impact on the composite properties. Thus, the aspect ratio of the platelets in the composite is an important tool which can be used
Fig. 15 a Representation of a computer model consisting of 50 randomly placed platelets. The platelets had an aspect ratio of 50 and filler volume fraction of 3% and b cross-section through the centre of the model. Reproduced from [30] with permission from Wiley
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Fig. 16 a Comparison between the experimental gas permeation values through epoxy and polyurethane nanocomposites as a function of filler volume fraction and the model predictions and b influence of misalignment on the permeation through the composites. Reproduced from [30] with permission from Wiley
to compare the different systems, different surface modifications, or different synthesis methodologies, etc. it can thus help to tune the microstructure development and to design the nanocomposites according to requirement. Determination of the aspect ratio is not possible from the transmission electron micrographs owing to significant extent of platelet bending, folding and misalignment. Other means to obtain average aspect ratio of the filler in the composite are thus required. Numerical finite element approach has been used to generate the computer models which can be used to generate information on the average aspect ratio of the platelets [30]. The computer models comprised of 50 non-overlapping identical platelets with aspect ratios of 50 or 100, which were randomly distributed and oriented in a periodic cubic box. The volume fraction of the platelets was varied between 1 and 5%. Figure 15a shows the representation a computer model consisting of 50 randomly oriented platelets with an aspect ratio of 50 and a filler volume fraction of 3%. Figure 15b also shows a cross-section through the center of the model. Figure 16a points towards the comparison of the experimental gas permeation values through epoxy and polyurethane nanocomposites as a function of filler volume fraction and the numerical predictions from the models as a function of the increasing misaligned filler platelets with different aspect ratios.
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The experimental and predicted values match very well till 3 vol% of the filler in the composite. At higher filler volume fractions, the comparison was observed to deteriorate owing to the lower extent of exfoliation experimentally. It was also observed that in this case, the filler in the polyurethane nanocomposites had higher aspect ratio than the epoxy nanocomposites resulting from the better interactions between the polymer and filler phases. Figure 16b also demonstrates the impact of misalignment on the permeation properties of the composites, the misaligned platelets were observed to be less resistant to permeation than the aligned platelets.
References 1. Mittal, V.: Barrier properties of composite materials. In: Mittal, V. (ed.) Barrier Properties of Polymer Clay Nanocomposites. Nova Science Publishers, New York (2009) 2. Osman, M.A., Mittal, V., Morbidelli, M., Suter, U.W.: Polyurethane adhesive nanocomposites as gas permeation barrier. Macromolecules 36, 9851–9858 (2003) 3. Xu, R., Manias, E., Snyder, A.J., Runt, J.: New biomedical poly(urethane urea)-layered silicate nanocomposites. Macromolecules 34, 337–339 (2001) 4. Chang, J.H., An, Y.U.: Nanocomposites of polyurethane with various organoclays: thermomechanical properties, morphology, and gas permeability. J. Polym. Sci. Polym. Chem. 40, 670–677 (2002) 5. Tortora, M., Gorrasi, G., Vittoria, V., Galli, G., Ritrovati, S., Chiellini, E.: Structural characterization and transport properties of organically modified montmorillonite/ polyurethane nanocomposites. Polymer 43, 6147–6157 (2002) 6. Wang, Z., Pinnavia, T.J.: Nanolayer reinforcement of elastomeric polyurethane. Chem. Mater. 10, 3769–3771 (1998) 7. Zilg, C., Thomann, R., Mulhaupt, R., Finter, J.: Polyurethane nanocomposites containing laminated anisotropic nanoparticles derived from organophilic layered silicates. Adv. Mater. 11, 49–52 (1999) 8. Chen, T.K., Tien, Y.I., Wie, K.H.: Synthesis and characterization of novel segmented polyurethane/clay nanocomposites. Polymer 41, 1345–1353 (2000) 9. Ma, J., Zhang, S., Qi, Z.: Synthesis and characterization of elastomeric polyurethane/clay nanocomposites. J. Appl. Polym. Sci. 82, 1444–1448 (2001) 10. Yao, K.J., Song, M., Hourston, D.J., Luo, D.Z.: Polymer/layered clay nanocomposites: 2 polyurethane nanocomposites. Polymer 43, 1017–1020 (2002) 11. Tien, Y.I., Wei, K.H.: Hydrogen bonding and mechanical properties in segmented montmorillonite/polyurethane nanocomposites of different hard segment ratios. Polymer 42, 3213–3221 (2001) 12. Tien, Y.I., Wei, K.H.: High-tensile-property layered silicates/polyurethane nanocomposites by using reactive silicates as pseudo chain extenders. Macromolecules 34, 9045–9052 (2001) 13. Tien, Y.I., Wei, K.H.: The effect of nano-sized silicate layers from montmorillonite on glass transition, dynamic mechanical, and thermal degradation properties of segmented polyurethane. J. Appl. Polym. Sci. 86, 1741–1748 (2002) 14. Osman, M.A., Mittal, V., Morbidelli, M., Suter, U.W.: Epoxy-layered silicate nanocomposites and their gas permeation properties. Macromolecules 37, 7250–7257 (2004) 15. Lan, T., Pinnavaia, T.J.: Clay-reinforced epoxy nanocomposites. Chem. Mater. 6, 2216–2219 (1994) 16. Messersmith, P.B., Giannelis, E.P.: Synthesis and characterization of layered silicate-epoxy nanocomposites. Chem. Mater. 6, 1719–1725 (1994) 17. Lan, T., Kaviratna, P.D., Pinnavaia, T.J.: Mechanism of clay tactoid exfoliation in epoxy-clay nanocomposites. Chem. Mater. 7, 2144–2150 (1995)
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18. Zilg, C., Mulhaupt, R., Finter, J.: Morphology and toughness/stiffness balance of nanocomposites based upon anhydride-cured epoxy resins and layered silicates. Macromol. Chem. Phys. 200, 661–670 (1999) 19. Brown, J.M., Curliss, D., Vaia, R.A.: Thermoset-layered silicate nanocomposites. Quaternary ammonium montmorillonite with primary diamine cured epoxies. Chem. Mater. 12, 3376– 3384 (2000) 20. Zerda, A.S., Lesser, A.J.: Intercalated clay nanocomposites: morphology, mechanics, and fracture behavior. J. Polym. Sci. Polym. Phys. 39, 1137–1146 (2001) 21. Kornmann, X., Lindberg, H., Berglund, L.A.: Synthesis of epoxy-clay nanocomposites: influence of the nature of the clay on structure. Polymer 42, 1303–1310 (2001) 22. Kornmann, X., Thomann, R., Mulhaupt, R., Finter, J., Berglund, L.: Synthesis of aminecured, epoxy-layered silicate nanocomposites: the influence of the silicate surface modification on the properties. J. Appl. Polym. Sci. 86, 2643–2652 (2002) 23. Kong, D., Park, C.E.: Real time exfoliation behavior of clay layers in epoxy-clay nanocomposites. Chem. Mater. 15, 419–424 (2003) 24. Ma, J., Yu, Z.Z., Zhang, Q.X., Xie, X.L., Mai, Y.W., Luck, I.: A novel method for preparation of disorderly exfoliated epoxy/clay nanocomposite. Chem. Mater. 16, 757–759 (2004) 25. Mittal, V.: Epoxy-vermiculite nanocomposites as gas permeation barrier. J. Compos. Mater. 42, 2829–2839 (2008) 26. Mittal, V.: Effect of the presence of excess ammonium ions on the clay surface on permeation properties of epoxy nanocomposites. J. Mater. Sci. 43, 4972–4978 (2008) 27. Osman, M.A., Atallah, A., Suter, U.W.: Influence of excessive filler coating on the tensile properties of LDPE–calcium carbonate composites. Polymer 45, 1177–1183 (2004) 28. Morgan, A.B., Harris, J.D.: Effects of organoclay Soxhlet extraction on mechanical properties, flammability properties and organoclay dispersion of polypropylene nanocomposites. Polymer 44, 2313–2320 (2003) 29. Kadar, F., Szazdi, L., Fekete, E., Pukanszky, B.: Surface characteristics of layered silicates: influence on the properties of clay/polymer nanocomposites. Langmuir 22, 7848–7854 (2006) 30. Osman, M.A., Mittal, V., Lusti, H.R.: The aspect ratio and gas permeation in polymer-layered silicate nanocomposites. Macromol. Rapid Commun. 25, 1145–1149 (2004)
Elastomeric Nanocomposites for Biomedical Applications Nicole Fong, Anne Simmons and Laura Poole-Warren
Abstract Elastomeric nanocomposites are gaining considerable attention as new materials for biomedical use. Elastomers such as polyesters, polyurethanes, and silicone rubber are excellent candidates as biomaterials in applications including tissue engineering due to properties such as their ease of synthesis and chemical manipulation, biodegradability and biocompatibility. However, when used alone these elastomers often fail to meet the mechanical and physical demands of the specific application. Elastomeric nanocomposites are composite materials comprising nano-sized reinforcements dispersed throughout the polymer matrix. The presence of such nanoparticles has been shown to improve the mechanical properties of the base elastomer as well as decrease their permeability properties making them more suitable for tissue engineering scaffolds and controlled drug release among other uses. Thus nanocomposite technology is creating greater applicability of elastomers for biomedical use, however continued research is required to better understand their behavior when material components are varied and the effect of nanoparticles on biological systems. Abbreviations NC MMT CNT QAC XRD TEM PHA
Nanocomposite Montmorillonite Carbon nanotube Quaternary ammonium compound X-ray diffraction Transmission electron microscopy Polyhydroxyalkanoate
N. Fong, A. Simmons and L. Poole-Warren (&) Graduate School of Biomedical Engineering, University of New South Wales, Sydney NSW, 2052, Australia e-mail:
[email protected] V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_10, Ó Springer-Verlag Berlin Heidelberg 2011
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PTMC UHMWPE PCL PGS FTIR TGA PHB HA MCNT PU CHX PDMS
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Poly(tetramethylene carbonate) Ultra high molecular weight polyethylene Poly(e-caprolactone) Poly(glycerol sebacate) Fourier transform infrared Thermogravimetric analysis Poly(3-hydroxybutyrate) Hydroxyapatite Multiwalled carbon nanotube Polyurethane Chlorhexidine diacetate Polydimethylsiloxane
1 Introduction Elastomers have been used in a range of biomedical applications since the 1940s as they have comparable properties to many tissues in the body, making them valuable materials in the tissue engineering, orthopaedics, and medical device fields. Elastomers are defined as polymeric materials with rubber-like properties, exhibiting an ability to elongate under applied force and return to its original state. Therefore, they offer the advantage of being easily formed into complex shapes and chemically manipulated to modulate properties such as compliance, tensile strength and stiffness [1]. In biomedical applications such as orthopaedics these capabilities make elastomers invaluable over the use of traditional metallic or ceramic materials. For example, elastomer–calcium phosphate composites have demonstrated more favorable elastic moduli to facilitate bone healing than their traditional metal counterparts. For metallic bone implants, stress-shielding is a common cause of bone refracture due to difference in elastic modulus between the bone and metal. In addition to the better mechanical compatibility, the elastomeric materials allow biological components such as bone marrow stromal cells, or mesenchymal stem cells to be incorporated into the implant to aid the healing process [2]. Despite these advantages, elastomers are often deemed inappropriate materials for certain biomedical applications due to limitations associated with poor diffusion or barrier properties of the pristine polymers and/or their long-term biocompatibility or mechanical resilience. For example, biomaterials currently used for urinary catheters are silicone, latex and silicone-coated latex. Latex catheters are easy and cheap to manufacture and they produce soft devices which are comfortable for the patient. The associated compromise is that a thicker wall size is required for sufficient rigidity. Latex is also associated with allergic reactions
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and is susceptible to infection and encrustation as well as having relatively poor tissue compatibility [3]. Notwithstanding these shortcomings, latex remains a common urinary device material due to patient comfort. The alternative, silicone, does not tend to induce allergic reactions and offers greater rigidity compared to latex, which allows for thinner walls and hence more efficient bladder drainage. On the downside, the higher rigidity can cause patient discomfort [4]. Nanocomposite (NC) technology has the potential to address such shortcomings of elastomers. The addition of nanoparticles within elastomers as well as other polymer matrices allows for a range of material properties to be modulated, including barrier, mechanical and thermal properties. This capability makes elastomeric NCs excellent candidates as biomaterials for medical applications since the desirable properties of the elastomer can be retained while other material characteristics may be tailored for the required application. Through careful selection of an appropriate polymer matrix, and selection and loading of nanofillers, mechanical properties that meet requirements of a particular application can theoretically be obtained. In this chapter, the most significant elastomeric biomaterials and their NCs studied to date are reviewed, including their impact on the material properties of the pristine polymer, and how nanotechnology is creating greater applicability of elastomers in biomedical applications.
2 Polymer Nanocomposites The origins of polymer NCs can be traced back to the early 1960s when Blumstein successfully adsorbed layers of methyl methacrylate onto the surface of montmorillonite (MMT) clay nanoparticles [5–7]. However, it was not until 1993 that the real potential of NCs was discovered and reported by the Toyota Central R&D Labs in Japan [8–10]. This group pioneered the synthesis of a polymer NC where homogeneous dispersion of MMT within a Nylon 6 (polycaprolactam) polymer matrix was achieved. The NC was reported to have significant improvements in mechanical and thermal properties at low loadings of 2–5% MMT, compared to pristine Nylon 6 and traditional Nylon 6 composites. At a loading of 4.7% MMT, tensile strength was seen to increase from 68.6 to 97.2 MPa, Young’s modulus from 1.11 to 1.87 GPa, flexural strength from 89.3 to 143 MPa, and heat distortion temperature from 65 to 152°C [11]. These findings attracted attention from numerous industries ranging from the automotive and aerospace to the packaging industries, and since then research into the development of polymer NCs, employing a variety of nanoparticles and matrix polymers for a large scope of applications, has gained momentum [12–17]. Polymer NCs are essentially comprised of two main components; (1) the polymer matrix material and (2) a filler component that has at least one dimension on the nanoscale [18], but in many cases a third organic modifier component is
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also required to distribute the nanofiller within the polymer matrix. A range of elastomeric matrices for NC systems have been studied including polyurethane, poly(e-caprolactone) and polysiloxane. Silicates are the usual nanofillers of choice in NC systems, although other fillers such as single and multi-walled carbon nanotubes (CNTs), alumina nanofibres, and carbon black have also been used. The silicates can be naturally occurring (e.g. montmorillonite) or synthetic (e.g. layered double hydroxides, fluorohectorite), provided at least one dimension, measures between 1 and 100 nm [19]. Clay minerals are classified according to their crystal structure and surface charge. MMT is one of the most frequently used clays for NC synthesis, and is considered to be a smectite belonging to the family of phyllosilicates, which are hydrous layered silicates of Al, Mg, Fe, and other elements, and large aspect ratios [18]. Smectites have a 2:1 layered structure made up of two tetrahedral sheets filled with silicon (Si4+) atoms edge-shared with a central octahedral sheet occupied by either aluminium or magnesium hydroxide. A single silicate layer is formed when the tetrahedral and octahedral sheets combine at the tips of the tetrahedrons of the silica sheets with hydroxyl layers of the octahedral sheets, making them in the order of 1 nm thick, with lateral dimensions ranging from 30 nm to several micrometres. Silicate nanolayers occur in nature in stacked configurations called tactoids. This, along with their hydrophilic nature makes them particularly incompatible and difficult to disperse in hydrophobic polymers. Therefore, organic modifiers are often introduced, whose role in polymer NC systems is to compatibilise the silicate surface with the polymer. As a result, dispersion of the silicate nanoparticles can be achieved. For this process, cationic surfactant molecules such as quaternary ammonium compounds (QACs, e.g. dodecyltrimethyl ammonium, hexadecyltrimethyl ammonium, dimethyl ditallow ammonium) are primarily chosen as their charged head groups can interact with the silicate particles, while their hydrophobic hydrocarbon tails facilitate silicate swelling and polymer intercalation (Fig. 1). Along with QACs, primary amines (e.g. dodecylamine, octadecylamine) have been the most widely used organic modifiers in NC fabrication.
Fig. 1 QACs such as octadecyltrimethyl ammonium chloride have a monocationic head group with a hydrophobic tail
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2.1 Polymer Nanocomposite Synthesis Several methods for polymer-NC synthesis have been described. These include an in situ polymerisation method, polymer melt intercalation method and solution cast method [20]. The in situ polymerisation method is a one-step method that was used in the pioneering studies by the Toyota group [8–10] in the synthesis of Nylon 6-MMT NCs. The monomer, in this case e-caprolactam, and nanoparticle, MMT, were combined prior to polymerisation to allow the monomer to intercalate between the silicate layers. The polymerisation process was then forced to occur by heating to 250°C, thus resulting in a polymer NC. Polymer melt intercalation involves heating of the polymer–clay mixture beyond the melting point of the polymer to enable the polymer chains to move to intercalate the clay layers. Vaia et al. [21] successfully demonstrated the use of this method to produce the first polystyrene-MMT NC. The solution cast method utilises an organic solvent to aid with the penetration of polymer chains between clay layers. Through the addition of the silicate and polymer to the solvent, the silicate layers expand and the dissolved polymer can intercalate between the layers. The subsequent removal of the solvent results in the polymer NC. Solution casting is commonly used for the manufacture of biomedical devices such as catheters, insulation for pacemaker leads and aortic balloons [22] and is a suitable method for these applications as it can produce high quality products at relatively low processing temperatures and cost. The low processing temperatures allow for the incorporation of bioactive agents, making them particularly attractive in medical applications. For NC synthesis, an additional advantage includes a relatively low cost increase per unit volume for significant improvements in material performance, particularly when abundant and naturally occurring nanoparticles such as silicates are used [18]. For these reasons, this method has been widely used by researchers in the NC field [23–26].
2.2 Dispersion Morphology Varying degrees of nanoparticle dispersion are usually achieved in polymer NCs. These are generally categorized as (1) conventional composite, (2) intercalated NC, or (3) exfoliated NC, as illustrated in Fig. 2. In conventional composites, the clay particles, for example, exist in agglomerated stacks or tactoids and there is no intercalation of polymer between clay layers, whereas intercalated NCs exhibit some insertion of polymer layers between the clay. Exfoliated NCs are achieved when the tactoids completely separate and the individual clay layers are dispersed throughout the polymer matrix in either a uniform or disordered manner with an average distance determined by clay loading [27]. To optimize the benefits that NCs have to offer with regards to material properties, the filler particles should be exfoliated throughout the matrix material.
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Fig. 2 Schematic diagrams of an (a) conventional composite, where nanoparticles remain agglomerated within the polymer matrix, (b) intercalated nanocomposite, where polymer chains intercalate between the nanoparticles, and (c) exfoliated nanocomposite, where tactoids completely separate and individual nanoparticles are dispersed throughout the polymer matrix in either a uniform or disordered manner with an average distance determined by filler loading
In an exfoliated form, maximal interfacial interactions with the polymer matrix can be attained due to the large aspect ratio of the clay, and therefore stress transfer to individual nanoparticles is increased, thereby contributing to the reinforcement and mechanical improvement of the materials. Exfoliation, however, is difficult to achieve in most polymer–clay systems since the clays’ preferred state is in stacked configurations, and being hydrophilic in nature they are intrinsically incompatible with hydrophobic polymers [28]. A major challenge in NC research then, is to determine effective ways in the manipulation and synthesis of these materials to optimize polymer intercalation between the clay layers.
2.3 Nanocomposite Characterization The characterization of polymer NCs for biomedical applications generally involves firstly the structural characterization of the materials followed by biological assays that are appropriate for the specific application. Structural analysis typically involves X-ray diffraction (XRD), transmission electron microscopy (TEM) and mechanical testing.
2.3.1 X-Ray Diffraction Through the analysis of XRD, or scatter, valuable information on the structure of polycrystalline materials can be obtained. In the case of NC characterization, XRD results give an indication of the dispersion morphology of the nanoparticles, and infer information on the orientation of the organic molecules in the intergallery space.
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XRD operates by directing an incident X-ray beam at the sample, which diverges from the X-ray source and passes through a divergence slit, to then strike and be diffracted by the sample. The diffracted beam passes through an antiscatter and receiving slit to be received by an X-ray detector. The divergence and antiscatter slits limit and direct the incident and diffracted beams and cause them to collimate, thus controlling the area of the sample that is exposed to the X-ray beam. The electronic detector converts the X-rays into pulses of electrical current which are processed to provide information about the intensity of the X-rays entering the detector [29]. Calculations used to determine the basal spacing of OMS or the average distance between silicate layers in NCs are based on Bragg’s Law (Eq. 1), where l is the known wavelength of the incident beam, 2h is the diffraction angle, or the angle between the diffracted beam and the transmitted beam, and d is the distance between silicate layers. l ¼ 2d sin h
ð1Þ
2.3.2 Transmission Electron Microscopy TEM provides a means to visually image the internal structure of the NCs. In polymer NC analysis, TEM is often used to complement XRD findings, as XRD gives a quantitative indication of average silicate dispersion and TEM can provide qualitative information on the dispersion morphology of the materials. TEMs enable imaging of specimens at very high magnification through irradiation by electron beams under vacuum. Depending on the electron densities within the sample material, the incident electrons either pass through the specimen (transmission electrons), or are scattered (scattered electrons). It is from these transmission and scattered electrons that high-magnification TEM images are formed to aid in the identification of NC morphologies [30]. Representative images of NCs obtained by TEM imaging are shown in Fig. 3.
Fig. 3 Representative TEM images of (left to right) elastomeric microcomposite, intercalated nanocomposite, partially exfoliated nanocomposite
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2.3.3 Mechanical Testing Testing of NC materials to assess their mechanical properties can be performed using a range of techniques including tensile and compression testing, hardness/ nanoindentation tests and tensile hysteresis and stress relaxation tests [31–34]. Of these, tensile tests are the most common and widely used method to gain an understanding of how a material will behave in a finished product. The results from the measurement of force as a specimen is subjected to deformation at a constant rate provide information on the ultimate tensile strength, elasticity and Young’s modulus of a material. The mechanical properties of NCs can give an indirect indication of silicate dispersion. Since the presence of the silicate layers within the material alters the mobility of the polymer chains, changes in their mechanical properties are observed [35]. In general the inclusion of nanofillers in polymer matrices causes UTS to increase, particularly at lower inclusion loadings, or to be maintained. Ultimate strain is seen to decrease in most polymer NC systems and Young’s modulus is enhanced. It should be noted that these findings are highly dependent on several factors including the matrix polymer properties and preparatory methods, which explains variations in results [36].
2.4 Elastomeric Nanocomposites as Biomaterials Among the aforementioned properties of elastomers that make them excellent for biomedical use, elastomeric NCs provide additional benefits particularly in terms of being able to tailor material properties for specific applications. As previously discussed, the mechanical properties of the NC materials can be modulated through careful selection of NC components and chemistry, and factors such as nanoparticle loading and method of synthesis [36]. Barrier properties are influenced in a similar manner [14, 36]. It is well understood that, with the addition of glass or mineral inclusions, an improvement in barrier properties is experienced in conventional polymer composite materials. However, other important material properties such as mechanical strength and thermal stability are usually sacrificed due to the weak interfacial interaction and stress concentrations that occur at the inclusion-polymer interface [37]. A key goal in the development of polymer NCs is to maintain or enhance the favorable diffusion properties of composite materials, as well as preserve other important properties that are sacrificed in traditional composites. In the attempt to achieve this ideal, the aspect ratio or dispersion of the inclusions is increased to encourage a higher degree of interfacial interactions and lower stress concentrations at the inclusion sites. In polymer NC systems, it is expected that a longer diffusive path is afforded by the presence of particles with increased aspect ratio, since the nanoparticle is assumed to be impermeable, thereby forcing the penetrant to traverse around them, illustrated in Fig. 4. Such improved barrier properties of
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Fig. 4 Enhanced barrier properties are achieved with better dispersion of the nanoparticle
NC materials make them particularly attractive for biomedical applications such as controlled drug delivery. Due to their favorable physical and chemical material properties, the most common elastomers studied for biomedical applications include polyesters, polyurethanes, and silicone rubber. The following section reviews current research into these elastomers and their NCs as biomaterials.
2.5 Polyester Elastomer Nanocomposites Polyesters commonly used for biomedical applications such as polyglycolide and polylactide, are limited in their use due to their stiffness and elastic deformation when exposed to long-term strain, as well as their acidic degradation products [38]. The development of thermoplastic polyester elastomers (TPEE), however, has helped to overcome these shortcomings whilst maintaining the favourable properties of the polyester. TPEEs are multiblock copolyether esters with alternating long or short-chain oxyalkylene glycols connected by ester linkages. Structurally, they contain repeating hard and soft segments much like polyurethanes and polyamide elastomers. The hard segments are typically composed of short-chain cyclic ester units such as tetramethylyene terephthalate and the soft segments are derived from aliphatic polyether glycols [39]. Due to the favourable low toxicity, appearance, and ease of processing and sterilizing of traditional polyesters, TPEEs with their additional improved biodegradation properties, elasticity and tensile strength, have gained considerable interest for use in biomedical applications such as controlled drug delivery, surgical sutures and tissue engineering scaffolds [38].
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An excellent example of the use of polyester elastomers in biomedical applications is in tissue engineering, where it is desirable for the tissue scaffold to have mechanical properties similar to that of the tissues they are proposed to replace, and be suitably compliant to accommodate physical and chemical deformations while being biocompatible [40]. Tissues have elastomeric properties and thus the development of elastomeric polymers for tissue engineering are of continuing interest, with the main biodegradable polyesters currently studied for these applications being polyhydroxyalkanoates (PHAs) [41, 42] and the more recently developed poly(trimethylene carbonate) (PTMC) [43, 44] copolymers [45–47], elastomeric poly(e-caprolactone) (PCL), poly(glycerol sebacate) (PGS) [48, 49] and poly(diol citrates) [50, 51]. Although elastomeric NCs are attracting much research interest for biomedical use, studies into NCs of these polyester elastomers are still relatively scarce. A selection of the few investigations into these elastomers as NC materials for biomedical use is discussed in this section. It is of note, however, that no current NC research using PGS has been reported in the literature to date.
2.5.1 Thermoplastic Polyester Elastomer PHAs have generated much research attention for applications such as controlled release, wound dressings and cardiovascular tissue engineering, as they are nontoxic biodegradable polyesters that may be easily chemically manipulated to modulate mechanical and degradation properties for a range of applications. Excellent reviews on PHAs and their applications are provided by Philip et al. [52] and Martin et al. [42]. PHAs are naturally produced from the cells of microorganisms and therefore have the advantage of being free from any toxic residues that may occur in the chemical synthesis of other polyesters, yet PHAs are still structurally comparable to chemically derived polyesters [42]. The development of PHA NCs for biomedical applications is relatively new, however it has been shown that the presence of dispersed nanoparticles within the PHA matrix affords improved material properties for medical use. Lim et al. [41] investigated an elastomeric polyester NC using the most commonly used PHA, poly(3-hydroxybutyrate) (PHB), and MMT nanoparticles to determine the effects of the nanoinclusions on polymer properties, in particular their thermal stability. NCs were solvent cast using PHB/MMT in the ratios 100/0, 97/3, 94/6 and 91/9 by weight and subsequently characterised using XRD, Fourier transform infrared (FTIR) spectroscopy and thermogravimetric analysis (TGA). Their results showed that intercalated PHB/MMT NCs were fabricated using low to moderate nanoparticle loadings, with increased thermal stability relative to pristine PHB. These results implicated the improved biostability of the materials, thus increasing their applicability as biodegradable biomaterials. PHA NCs with nano-hydroxyapatite (HA) reinforcements have also been studied for more specific orthopedic applications. Chen et al. [53] investigated the
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mechanical properties and bioactivity of a solution cast NC based on bioresorbable polymer-poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBHV) and nano-sized hydroxyapatite (HA) for bone replacement. Structural characterization of the materials were conducted as well as thermal and dynamic mechanical tests. Results from these studies showed that the HA nanoparticles were well dispersed within the PHBHC matrix, which was hypothesized to be responsible for the increased degradation temperature, and thus improved biostability, of the materials as well as the increased storage modulus. Bioactivity of the PHBHV-HA NCs were assessed by immersing materials into a simulated body fluid with solution ion concentrations similar to those of blood plasma as observing the formation of surface calcium-phosphate. A thin layer of calcium-phosphate was observed on the NC surface after 30 days, indicating good bioactivity.
2.5.2 Poly(trimethylene carbonate) Poly(trimethylene carbonate) (PTMC) has been studied as biodegradable surgical sutures [45], controlled drug delivery vehicles [46, 47], bone cement [54], as scleral buckles [55] and as tissue engineering scaffolds [56]. Among studies into PTMC copolymers as tissue engineering scaffolds, Pego et al. [57, 58] demonstrated that these elastomers show potential in their use in long-term degradable devices for soft tissue engineering as they are both biocompatible [57] and mechanically resilient [58]. Ribeiro et al. [59] developed a NC material based on PTMC reinforced with HA and multiwalled carbon nanotubes (MCNT) for joint repair, where HA was used to stimulate osteoinduction and the MCNT was used as a reinforcing agent and to provide mechanical damping. The nanoparticles were incorporated into the polymer matrix by in situ polymerization and materials were assessed for their coefficient of friction and surface properties relative to cartilage and the conventionally used biomaterial for joint repair, ultra high molecular weight polyethylene (UHMWPE). Results suggested that the friction coefficient of PTMC composite materials were closer to that of cartilage than UHMWPE and surface behavior of the NCs was also comparable to articular cartilage. However, surface attraction forces of UHMWPE and the PTMC NCs were found to be lower than that of natural cartilage, implicating poor lubrication properties.
2.5.3 Elastomeric Poly(e-Caprolactone) Amsden and colleagues [43, 44] synthesised and characterised biodegradable thermoset elastomers using a star copolymer of D,L-lactide and e-caprolactone as potential biomaterials for drug delivery and tissue engineering applications. They found that the elastomers’ mechanical and degradation properties could be modulated through different crosslinking densities, and materials with properties
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comparable to elastin could be produced making them excellent candidates for soft tissue engineering applications. Di et al. [60] also demonstrated the possibility of modulating the mechanical and barrier properties of PCL matrix through incorporation of organically modified silicate nanoparticle. In their study Young’s modulus and ultimate strength increased with increased nanoparticle loading and elongation at break decreased and the materials’ permeability to air was observed to significantly decrease. In further assessment of the material properties of PCL NCs in more specific biomedical applications, Kim and Fuchs et al. investigated their biodegradation and biocompatibility properties for bone regeneration. Kim [61] synthesized and characterized exfoliated HA/PCL NCs where oleic acid was used as the organic modifier. Morphological features along with the mechanical properties and cellular responses compared to traditional HA/PCL composites were examined with results showing that the NCs had better mechanical properties compared to the traditional composites and pristine PCL as well as improved cellular interactions, where the oleic acid was thought to improve osteoblastic response. Fuchs et al. [62] assessed the suitability and biocompatibility of PCL NCs with calcium-deficient hydroxyapatite (CDHA) nanoparticles in 11 and 24% PCL matrices for bone tissue engineering and vascularisation through investigating their ability to support the cell growth of endothelial cells and primary osteoblasts under monoculture conditions. Biocompatibility tests were conducted using human umbilical vein endothelial cells (HUVEC) and primary osteoblasts, with results showing that the HUVECs on the 24% PCL NCs formed a more homogeneous endothelial monolayer the materials with 11% PCL content. Osteoblasts remained viable on all material surfaces, with no clear effect of material composition. Further, assessment of the formation of pre-vascular structures using co-cultures of the endothelial cells and osteoblasts suggested that this was time and PCL-loading dependent as more pre-vascular structures were observed after 1 week on the 24% PCL NC materials, whereas after 4 weeks of pre-vascular structure formation was slightly higher on the 11% PCL materials.
2.5.4 Poly(Diol Citrates) Poly(diol citrates) are a biodegradable elastomer recently developed by Yang and colleagues [48] to specifically address mechanical requirements of scaffolds for tissue engineering, particularly for small diameter blood vessels. The elastomer, poly(1,8-octanediol-co-citric acid) (POC) is non-toxic, and easily synthesised with mechanical and barrier properties that maybe be modulated. In the characterization of the POC elastomer, Yang et al. [49] found that the mechanical properties of the materials could be modified through the selection of diols and the postpolymerization conditions. However it was also determined that the mechanical properties of the POC elastomer may not be mechanically sufficient as scaffolds for tissues that experience significant tensile or compressive forces such as cartilage. To enhance the mechanical properties of the elastomeric poly(diol citrate) Webb et al.
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[63] used nanocomposite technology, incorporating a nanofibrous poly(L-lactic acid) (PLLA) nanophase to the POC matrix. Results from tensile tests showed an increase in tensile strength, modulus, and elongation at break mechanical properties comparable to those of human cartilage, ligaments, and blood vessels. Additional assessment of the compressive modulus of the NC materials was found to be similar to human and bovine articular cartilage. Thus these studies demonstrated that poly(diol citrate)–PLLA NC scaffolds hold much promise for tissue engineering.
2.6 Polyurethane Nanocomposites Thermoplastic polyurethanes (PUs), are considered to be one of the most biocompatible and haemocompatible materials currently available, and their excellent physical–mechanical properties such as toughness, fatigue resistance and durability, and chemical versatility that allows for the incorporation of bioactive molecules, also make them popular choices in the manufacture of biomedical devices such as the left ventricular assist device, endotracheal tubes, catheters and vascular grafts. PUs are defined by the presence of urethane linkages and their properties can range from soft elastomeric to hard thermosetting materials. In general, biomedical PUs have a segmented copolymer chemistry made from a combination of a diisocyanate, macrodiol and a chain extender. Figure 5 shows their characteristic alternating polydisperse ‘soft’ and ‘hard’ microdomains, which are attributed to the macrodiol, and the diisocyanate linked with the chain extender, respectively. Since both segments are chemically incompatible, a phase separated morphology results, with the ‘soft’ macrodiol phase conferring elasticity, lubricity and softness, and the ‘hard’ domains providing cohesive strength to the material [64]. The first biomedical PU was presented by Boretos and Pierce [65] who found that a segmented PU experienced no failure or change in tensile properties after 9 days of implantation in calves. They also discovered that the material remained essentially free from thrombus formation and emboli when used as atrial cannulas Fig. 5 Characteristic alternating polydisperse ‘soft’ and ‘hard’ microdomains of polyurethane, which are attributed by the macrodiol, and the diisocyanate linked with the chain extender, respectively
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and vascular tubing in calves for 1 week and also when used in the vena cava of dogs. Subsequent experiments also showed no tumor induction in response to the material. A later study by Boretos et al. [66] further confirmed the use of PU as potential biomaterials, performing particularly well in cardiovascular applications. Since then, the study of PU for biomedical applications has confirmed that these materials are one of the best available elastomers for biomedical applications [67–69] even though they have been found to degrade in vivo via hydrolysis, oxidation and stress-cracking [70, 71]. PU NCs have been studied for a range of biomedical applications including infection control [72–74] and controlled drug delivery [75, 76]. Poole-Warren et al. [72] demonstrated the use of bioactive agents as organic modifiers of MMT nanoparticles within a poly(ether urethane) matrix to impart biological properties to the polymer matrix as well as aid dispersion of the MMT throughout. More specifically, the antibacterial agent chlorhexidine diacetate (CHX) and the non-antibacterial compound dodecylamine (12CH3) were tested against the pathogen Staphylococcus epidermidis, and in platelet and cell adhesion studies respectively. For both organic modifiers, NCs were solvent cast and were found to have intercalated to exfoliated structures. PU-CHX-MMT NCs showed an excellent 2-log reduction in adherent bacteria compared to control PU materials, and PU-12CH3-MMT materials demonstrated a decrease in platelet and fibroblast adhesion. It was therefore shown that organic modifiers with appropriate structure may be used to alter biological interactions of PU NCs. Fong et al. [73] expanded on this research and reported on the use of PU NCs as antibacterial biomaterials for medical devices such as urinary catheters. It was shown that molecules with appropriate structure may be dual-functional within a NC system, behaving as both a silicate dispersant and bioactive agent. More specifically, Fong et al. showed that the antibacterial dicationic molecule, chlorhexidine diacetate, produced intercalated to partially exfoliated NCs when used as an organic modifier and/or additive, and promoted a 2-log drog in Staphylococcus epidermidis, an opportunistic microorganism significant in medical device-related infections [77]. Other similar studies have been reported supporting such findings, including a study by Styan et al. [74] who used the quaternary ammonium compound, EthoquadÒ O/12PG, with an 18 carbon alkyl chain as the organic modifier. Another approach to conferring antibacterial properties to PU NCs is through the addition of nanoparticles with antibacterial properties. Hung and Hsu [68] demonstrated this technique using silver nanoparticles as the filler component in a poly(ether urethane) matrix. Their assessment of the biocompatibility and antibacterial properties of the solvent cast NC films suggested that the presence of silver nanoparticles within a PU matrix improved the elastomer’s biocompatibility and cell proliferation properties, and reduced macrophage attachment, which may improve material biostability [68]. Permeability of biomaterials to water and air can potentially lead to problems in implanted devices. Xu et al. [78] demonstrated the advantages of nanocomposite technology in significantly decreasing the materials’ permeability characteristics.
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NCs were prepared using poly(urethane urea)s (PUU) and two organically modified layered silicates, Cloisite 15A (Southern Clay Products) and Nanomer I.30 TC, which use alkyl ammonium compounds dimethyl ditallow ammonium and octadecylammonium as organic modifiers, respectively. Structural assessment of the solvent cast NCs suggested that intercalated and exfoliated NCs were formed and Young’s modulus could be modulated with silicate loading while maintaining the tensile strength and elasticity of pristine PUU. Permeability studies showed a five-fold reduction in water vapour permeability at the highest silicate loading (20 wt%) compared to the unfilled PUU control materials. Williams et al. [79] studied the potential of PUNCs incorporating multiwalled carbon nanotubes (CNTs) for electrical interfacing with neural tissue. Traditional platinum electrodes used for this application are associated with fibrous encapsulation which increases the impedance at the neural interface. In the study by Williams et al., NC technology was hypothesised to afford conductive properties to the biocompatible PU matrix while maintaining the processibility and other favourable properties of the PU elastomer. Although homogenous dispersion of the CNTs throughout the PU matrix was not confirmed, minimal cell growth inhibition of L929 fibroblasts was observed relative to pristine PU and the conductivity of films was found to range from being non-conductive at 0% loading of CNTs to 6.55 9 10-2 S/cm at 20 wt% loading suggesting that these NC materials show potential as electrically conductive biomaterials.
2.7 Silicone Nanocomposites Silicones are synthetic polymers with a backbone of repeating silicon and oxygen bonds, referred to as siloxane. Silicon is also typically bound to methyl groups, or other organic groups. The presence of such organic groups to the inorganic polymer backbone equips silicone with its unique properties such as thermal stability and biocompatibility [80]. The most commonly used silicone is polydimethylsiloxane (PDMS), which are linear polymers with a repeating unit of –(Si(CH3)2)O)–. These polymers can be transformed into elastomeric silicone structures through cross-linking reactions with radicals, by condensation or addition-cure reaction. Details of elastomeric silicone synthesis have been described elsewhere [80]. Silicone rubber has been the most frequently used and widely studied polymeric biomaterials, for applications including controlled drug release [81], bone bonding and tissue engineering scaffolds [82] and conductive electrodes for neural stimulation [83]. It is considered to be an excellent elastomer for biomedical purposes as it is known to be biocompatible with excellent biostability [84] and possess good elastomeric properties over a wide temperature range, thermal stability, excellent resistance to biodegradation, oxidation and ageing [1]. Despite these properties, silicone is associated with poor bioactivity and mechanical properties that are often insufficient for certain applications. NC technology allows these
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properties to be modulated through the addition of nano-reinforcements and bioactive agents. Meng et al. [85, 86] demonstrated the biofunctional capabilities of PDMS NCs by incorporating bioactive agents in the organic modification process. More specifically, PDMS/MMT NCs were synthesized using different antifungal agents as organic modifiers in the development of NC materials for antibacterial urinary catheter coatings. In their study using a topical and active allylamine antifungal agent, terbinafine hydrochloride, as the organic modifier of MMT nanoparticles, exfoliated NCs with enhanced thermal stability and excellent antifungal properties against Candida albicans were formed using an in situ polymerization technique [85]. In another study by Meng et al. [86], the antimicrobial agent chlorhexidine acetate was used as the organic modifier. Exfoliated NCs were again successfully formed using similar processing methods and an antibacterial effect was reported for the NC materials against Staphylococcus aureus and Escherichia coli, whereas control PDMS materials did not demonstrated any antibacterial activity. Mechanical properties of the NC materials were improved in the presence of the nanoparticles in both studies, increasing tensile strength and elongation at break by up to 232, 370%, respectively. Han et al. [87] aimed to enhance the bioactivity and bone-bonding properties of octamethyltrisiloxane for bone regeneration by incorporating HA nanoparticles into the polymer matrix. Structural characterization of the NCs showed that the HA nanoparticles were uniformly distributed throughout the silicone rubber matrix and mechanical tests indicated an improvement in tensile strength and Young’s modulus as well as maintaining the high elongation of the silicone elastomer on the addition of HA nanoparticles. In terms of bioactivity the interaction of preosteoblasts with NCs was assessed. The presence of attached osteoblasts confirmed the biocompatibility of the control silicone and NC materials, and cell density on the NC materials significantly increased with time relative to pristine silicone, thus indicating the superior biological properties of the silicone incorporating HA. Depan et al. [88] combined the favorable biodegradable, biocompatible, nontoxic, and biofunctional material properties of the naturally occurring polymer, chitosan (CS), with the excellent properties of PDMS in a CS-PDMS-MMT NC to overcome the limitations of using CS alone in wound dressing. These limitations including its solubility only in aqueous acidic solutions and rigid structure could theoretically be modulated using NC technology, leading to a material with greater applicability as wound dressings. NC materials were made by the intercalation of CS between MMT layers, followed by UV-irradiated grafting of PDMS onto CS. Structural characterization of the NC materials showed that an intercalated NC was formed with increased thermal stability and hydrophilicity as well as water absorption properties that could be modulated by MMT loading. Zhou et al. [89] fabricated a PDMS NC incorporating silver-chitosan -clay nanoparticles to investigate the material’s antibacterial effect on bacteria commonly associated with catheter-associated infections; Escherichia coli, Pseudomonas aeruginosa, Staphylococcus aureus and Candida albicans. Exfoliated NCs were made using a clay to silver-chitosan ratio of 1:1.3 through
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crosslinking at room temperature under vacuum. Excellent antibacterial activity of the NC materials against E. Coli, P. aeruginosa and S. aureus was observed compared to AgNO3 and studies of Ag release into artificial urine indicated retarded rate of release in NC material relative to PDMS-silver-chitosan materials in the absence of clay. In another study by Zhou et al. [90], the PDMS film was reinforced with MMT organically modified using phosphatidyl choline (PC), a phospholipid in eukaryotes responsible for the fluidity and thickness of eukaryotic membranes. Considering the origins of PC, its haemocompatible properties make it an excellent candidate for use in antithrombogenic applications. NCs were made by mixing PDMS and the PC-MMT nanoparticles in solvent and crosslinking under vacuum at room temperature. The PC-MMT nanoparticles were believed to be well dispersed within the PDMS matrix, and mechanical properties of PDMS were maintained in the NCs. Haemolysis tests to assess blood compatibility of the NCs demonstrated lower haemolysis in these materials compared to pristine PDMS, thus suggesting potential for the PDMS-PC-MMT NCs as antithrombotic biomaterials. Another biomedical application in which NC technology may prove to be beneficial is in transdermal drug delivery using pressure sensitive adhesive (PSA) systems. Currently, this means of drug delivery is associated with problems with drug permeation through the protective stratum corneum layer of the skin, and poor adhesive properties. Shaikh et al. [76] studied the role of a PDMS based NC as a PSA system for transdermal drug delivery. MMT nanoparticles were prepared using octadecylamine as an organic modifier, and NCs were subsequently fabricated using a range of nanoparticle loadings (0, 2, 5, and 10 wt%) by solvent casting. It was hypothesized that the NC materials would offer increased control over the drug release kinetics and adhesive properties relative to unfilled PDMS based PSA materials. Results showed that partially exfoliated NCs were achieved, promoting a more uniform release rate and significantly reduced the initial burst release from the materials compared to pristine PDMS. Further, assessment of the NC materials’ shear strength showed more than 200% improvement at the lower clay loadings, implicating improved adhesive properties.
2.8 Styrenic Elastomer Nanocomposites Styrenic elastomers are block copolymers containing polystyrene or polyethylene hard segments, and butadiene, isoprene, butylene or ethylene butylene soft segments. In 1969, the Shell Oil Company developed the styrenic block copolymer, Kraton, based on isoprene and butadiene [91]. It was found that these materials were not feasible as biomaterials due to the poor oxidative stability of these soft segments. The subsequent introduction of an ethylene–butylene segment to the copolymers saw an improvement in oxidative stability, leading to the development of poly (styrene-b-ethylene butylene-b-styrene) (SEBS) block copolymer. As well
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as SEBS, Poly(styrene-b-isobutylene-b-styrene) (SIBS) block copolymers are among the most commonly studied styrenic elastomers. SIBS was first reported by Kennedy in 1989 [92] and has attracted attention in the biomedical field due to its biostability and biocompatibility [93–96]. The use of SIBS as a biomaterial was first studied by Pinchuk et al. [97] They reported that the materials possessed superior biostability due to its quaternary carbon backbone of the polyisobutylene polymer. Since then, they have be used in a range of biomedical applications including drug eluting coatings for stents [98]. Further details of SIBs as biomaterials can be found in an excellent review by Renade et al. [99]. Research into styrenic elastomer NCs is still relatively scarce, however studies that have been conducted on SEBs based NCs have demonstrated similar improvements to material properties as other polymer-NC systems. More specifically, Hasegawa and Usuki [100] reported on SEBS-MMT NCs made using a compression moulding technique. The resulting NCs were characterized using TEM and XRD. Results showed that the polystyrene microsegments of the SEBS polymer in the NC materials were arranged along the length of the nanoparticle inclusions, suggesting a level of control of material structure and properties could be achieved. In terms of structural and mechanical properties, Ganguly et al. [101] studied SEBS-MMT NCs incorporating a series of organically modified and commercially available clay nanoparticles, Cloisite 10A and Cloisite 20A. Results from mechanical and structural characterization studies showed that exfoliated NCs could be achieved in the SEBS at lower loadings of Cloisite 20A nanoparticles and significantly improved dynamic mechanical properties as well as superior Young’s modulus and tensile strength was achieved using Cloisite 20A compared to the other nanoparticles investigated and unfilled SEBS. To the best of our knowledge, SIBS based NCs have not yet been reported on in the literature, however given the excellent biostability and biocompatibility properties of these styrenic polymers, research into SIBS-NCs as biomaterials will certainly flourish in the near future.
3 Limitations and Future Directions Current research in elastomeric NCs for biomedical applications shows great promise in their use in clinical settings due to the favorable material properties that may be achieved through nanoparticle reinforcement of the pristine polymer. However inconsistencies in their success for different applications need to be better understood. This level of understanding is well under way although the difficulty in this task results from the endless combinations of elastomer chemistry, nanoparticle type and synthesis technique as well as the specific application for which the material is destined to be used. In terms of biocompatibility, many studies have been concerned with the toxicity of the NC materials intended for biomedical use and the degradation products
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of their polymers [102–104], however problems facing NCs for biomedical use is now concerned with how nanoparticles interface with biological systems. Studies into nanotoxicology are still in their infancy although research to date highlights the pressing need to understand the fate of nanoparticles in different biological systems and applications [105, 106]. In a review article by Lanone and Boczkowski [105], the properties of nanomaterials suggesting potential nanotoxicology issues in biomedical applications were discussed. In general, the factors influencing the potential toxicology of nanoparticles include their chemical composition and surface reactivity, nanoparticle size and shape, and particle retention time [105, 107]. Of particular interest in current nanotoxicology research is CNTs. In a review paper by Smart et al. [107] research into the toxicity and biocompatibility of CNTs in biomedical applications was discussed. Overall, although no nanotoxicity of CNTs has been confirmed to date, in vivo studies have shown adverse effects of CNTs on the body. For example CNT lung inflammation and granuloma formation was observed when CNTs were introduced into the trachea of mice and rats [108–110]. Results from such studies warrant continued investigation into the fate of nanoparticles in biological systems to be able to properly understand the behavior and limitations of NC materials for biomedical applications. Thus, in addition to the excellent properties of elastomeric NCs that have been evident in the literature to date, results from a more detailed understanding of the behavior of nanoparticles within the human body will provide a more complete understanding of the materials, potentiating a more confident and widespread use of elastomeric NCs as biomaterials for a range of applications.
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Actuators and Energy Harvesters Based on Electrostrictive Elastomeric Nanocomposites Kaori Yuse, Pierre-Jean Cottinet and Daniel Guyomar
Abstract Research and development efforts devoted to electro active polymers (EAPs) are being actively undertaken today due to the numerous advantages of these materials. Moreover, from the viewpoint of world-wide ecological tendencies, renewable and clean energy sources turn the heads of not only researchers. The harvesting or scavenging of ambient energy constitutes an important alternative stage. Among various specifics of EAPs, the point of a high strain level is being considered in order to render EAPs promising materials for actuators or energy harvesters. Piezo devices are usually used for such applications, but when comparing their strain level to that of EAPs, i.e., around 0.2% for piezo elements and easily more than 20% for EAPs, the replacement is highly interesting. Although the Young modulus is smaller than that of ferroelectric materials, the potential storage energy remains higher. EAPs thus show much promise, but so far, the application of high electrical voltages is required and the utilization with other electrical components is limited. To overcome these drawbacks, an intermediate material was developed, exhibiting a large strain at reasonable levels of applied voltage. An investigation of the power harvesting with EAPs was also carried out. Keywords Electroactive polymer (EAP) Dielectric polymer Conductive filler Nano-carbon Low-powered actuator
K. Yuse (&), P.-J. Cottinet and D. Guyomar Université de Lyon, INSA-LGEF, 8 rue de la Physique, 69621 Villeurbanne, France e-mail:
[email protected] V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_11, Ó Springer-Verlag Berlin Heidelberg 2011
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1 Introduction 1.1 EAPs as Actuator Materials Today, there are several active studies on polymers capable of being actuated by induction of electrical fields. These polymers are called electro active polymers (EAPs), and their most interesting characteristic is their ability to be largely deformed. EAPs can achieve strains up to several hundred percent. They are thus of interest to be used as sensors/actuators where large deformations are necessary. As sensor/actuator materials, many kinds of transitions are possible, and principal types include mechanical to electric, magnetic or thermal energy. These transitions can be obtained by a large variety of materials, e.g., piezo-elements, optical fibers, or shape memory alloys (SMAs). They can be used in a wide range of applications. As materials for the transition between mechanical and electrical energy, the piezo element is the most frequent and widely commercialized material due to its numerous advantages, such as its high precision, high electro-mechanical responses, and its elevated response speeds. PZT (Pb(Zrx, Ti1-x)O3) appeared in the 1950s. This is a ceramic material that can obtain much bigger deformations than those of the already existing piezo element, BaTiO3. Soon after its appearance, it took the place as main sensor/actuator material. Many other components have also been investigated. With the ecological tendencies of today, lead-free piezoelectric materials are in demand. Their deformation level is, however, currently not high enough to replace the ceramics comprising lead. Polymer materials, such as PVDF, also exist. Nevertheless, their low strain level around 0.2% is regrettable. With a crystal structure, the strain level can be made to increase but no further than to approx. 1.7%. For applications of micro-structures, PVDF is an ideal material but not for applications of, for example, artificial human muscles where an extra large strain is required. The shape memory alloy (SMA) is famous for being a large-displacement sensor/actuator material. Still, its maximum strain is limited to around 8%, and because of its rapid fatigue, the expected strain is generally lower than this when the material is employed under cyclic loading. Despite this, SMA has another specific advantage: it can generate a large force. The response speed is quite small but it can be salved by miniaturelization. Nonetheless, the difficulty of temperature control and its large energy requirement need to be improved or solved in order for it to be discussed for the same application fields as PZT. EAPs then appeared which can generate extra large deformations. The studies on EAPs started mainly in robotics and especially for the medical field. The human muscle can generate around 20% of deformation. An air pump method has been often used for the replacement but a smooth movement was difficult to create. EAPs can realize such smooth and flexible actuation with the strain level easily exceeding 20%. The advantage of EAPs is not only their amazing strain level:
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they are also light-weight, inexpensive, with a convenient power source, fracture tolerance, softness, and ease of processing and forming. A small margin compensation is another weak point for ceramic piezo elements but EAPs can efface this factor. The stored potential energy is described by YS2/2. The Young modulus Y of piezo element is significant, 90–100 GPa (PZT), but with its low strain S level, 0.2%, the energy ends very small. Instead, the Young modulus of the polymer is generally low, up to some hundreds MPa by PU [1]. However, its large strain compensates for the low energy. Compared to shape-memory alloys (SMAs), EAPs demonstrate a higher response speed, lower density, and easier control mechanism. Instead of magnetic energy transition, the loss is quite low. However, the scope of practical applications of EAPs is limited by a low actuation force, a low mechanical energy density, and a low robustness. Progress toward actuators being used in robotic applications with performances comparable to those of biological systems will lead to great benefits [2]. The application field expands day by day, and there are many domains which have awaited such a promising large-deformation actuator/sensor material.
1.2 Energy Harvesting and EAP This chapter presents the use of EAPs as energy harvesting materials. The concept of energy harvesting is briefly introduced here. The energy issue has always been a potential threat to society ever since industrialization. During the past 10 years, much research has been performed on renewable and clean energy sources in order to gradually eliminate fossil fuels from megawatt power plants. While megawatt power plants generate electricity with increasing dimensions, the proliferation of portable or wearable devices are going toward the other end of the scale. Moore forecasted a doubling in the number of transistors in a given area of silicon approximately every 2 years. This prediction, now known as Moore’s Law, has held true for approximately four decades. As Moore’s Law indicates, a brand new generation of digital devices has followed the improvements in semiconductor technology. Small size scales, compact storage densities, minimized energy consumptions and shorter processing times are all direct results of that evolution. Consequently, the recent progresses in ultra low-power electronics allow the powering of complex systems using either batteries or environmental energy harvesting. The harvesting or scavenging of ambient energy provides an important alternative for the paradigm of primary electricity generation strategies and storagebased supply. Significant research efforts and industrial development have led to energy harvesting based on piezoelectric materials. This is one of the most promising solutions for direct power supply and energy storage for low-power wearable devices. Initially, the electrical stimulation of polymers produced relatively small strains or harvesting energies, restricting their practical use. But nowadays, polymers
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exhibiting large strains have been elaborated and show great potential and capabilities for the development of practical applications. Active polymers which respond to electric stimuli, i.e., EAPs, exhibit deformations two to three orders of magnitude higher than the striction-limited, rigid and fragile electroactive ceramics (EACs). From a technological point of view, the generator mode of electronic EAPs is potentially as important as the actuator mode. Actuators are indeed pervasive in modern technology, yet the critical need for new energy systems, such as generators, may be of larger significance than the sheer number of possible applications.
2 Variety and Principles of EAP Actuators Mainly, EAPs used for harvesting energy belong to one of the following two types: (1) dielectric elastomers and (2) electrostrictive polymers. This section explains their principles one by one, which have been thoroughly reviewed in the cited references. Also, some of the most recent developments for certain polymers are presented, and various applications of active polymers are given.
2.1 Dielectric Elastomer Figure 1 illustrates the generator mode of dielectric elastomers. The dielectric elastomer is stretched by an outside mechanical force, as shown in Fig. 1a. In the stretched state, opposite charges are induced on the opposing compliant electrodes using a voltage difference. The dielectric elastomer is now allowed to contract using elastic restoring forces. Note that in the contracted state, shown in Fig. 1b, the dielectric elastomer is thicker and has a smaller area relative to the stretched state. Assuming that the charge is constant during the contraction process, the change in geometry has caused a separation of opposite charges by a greater distance (the increase in thickness by contraction), and has compressed like charges into a smaller area (the decrease in area by contraction). Both changes increased the electrical energy, and hence the stored electrical energy. The
Fig. 1 The principle of dielectric elastomer generators: a stretched (charge at low voltage) and b contracted (charge at high voltage)
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mechanical work has been converted to a greater amount of stored electrical energy, and the material has operated as a generator transducer. The qualitative physics analysis of the generator mode, described above, can be analyzed electrically in quantitative terms using conventional macroscopic lumped parameters. Although electrical lumped parameters such as capacitance are more easily manipulated in the engineering sense, the physics of the dielectric elastomer generator mode can provide greater insights into more complex phenomena, some of which are described below. Parameter models of dielectric elastomer generators typically start, as they do for the actuator mode, by considering the capacitance of the electroded film. The capacitance in the stretched or contracted (relaxed) states, can be written as C¼
e 0 er A t
ð1Þ
where e0 is the permittivity of free space, er is the relative permittivity of the material, A is the stretched or contracted area where the opposite electrodes overlap (called the ‘active area’), and t is the polymer thickness. The electrical energy of a capacitor We with charge Q can be written according to the following well-known formula; We ¼
1 Q2 1 t Q2 ¼ 2 C 2 ðe0 er AÞ
ð2Þ
The polymer can be described by the constant volume approximation [3]; i.e., A t ¼ Vol ¼ constant
ð3Þ
where Vol is the volume of the elastomer. The change in electrical energy dWe, due to an incremental change in state can be then be written as dWe ¼
Q We dA ðconstant volumeÞ dQ 2 A C
ð4Þ
dWe ¼
Q We dt ðconstant volumeÞ dQ þ 2 t C
ð5Þ
or equivalently,
The term (Q/C)dQ is merely the incremental change in the electrical energy due to charge flowing onto the film at voltage V = Q/C. This term also includes any leakage of charge through the film. Equations 4 and 5 are very general, given the constant volume assumption, but further insight into the generator behavior and the important parameters can be obtained by examining simplified cases. Consider a constant charge system with charge Q. Using Eq. 2 directly for the electrical energy stored in the film, the difference in electrical energy between stretched and contracted states, i.e., the generated energy Wg, can be written as
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Wg ¼ Wcontracted Wstreched ¼
1 2 1 1 Q 2 Ccontracted Cstreched
ð6Þ
Equation 6 can be rewritten as Wg ¼
1 Q2 Cstreched Cstreched 1 ¼ Wstreched 1 Ccontracted Ccontracted 2 Cstreched
ð7Þ
Summarizing the basics of the dielectric elastomer generator mode, an initial charge is placed on a dielectric elastomer film that is stretched with regard to its area [4]. The film is then allowed to contract, which in turn further separates opposite charges and compresses like charges to a smaller area. From a macroscopic lumped parameter perspective, contraction reduces the capacitance. Using either viewpoint, if the charge is constant on the system, the contraction causes an increase in the stored electrical energy by raising the voltage of the charge. This increase in electrical energy is due to the fact that mechanical forces work against the electric field pressure. With a suitable electrical circuitry, the increase in stored electrical energy can be harvested to produce a net generated electrical energy output from the system. It is implied in this description that the system typically cycles. In other words, the film is stretched again by mechanical forces, an initial charge and energy is placed on the film, and this is followed by a new round of contraction. Dielectric elastomer leakage is also a practical consideration in power generation. Leakage is a direct loss to the system. The importance of leakage phenomena depends very much on the dielectric elastomer resistivity and the frequency operation. Lower frequency operation requires lower leakage losses and higher polymer resistivities. This is because lower frequencies allow more time for leakage losses to accumulate relative to the energy-producing cycle. Leakage losses may influence the choice of polymer, but a good dielectric elastomer can generally be identified to address this issue for most applications. Perhaps the biggest advantage of dielectric elastomers is the point that, in addition to a good dynamic range, they operate best at relatively long strokes and modest forces. It is difficult to address this part of the design space using conventional smart materials such as piezoelectrics. Moreover, it is a very common mechanical input available from a number of sources in the environment, such as human motion and wave power.
2.2 Electrostrictive Polymers Piezoelectric ceramics (such as PZT) have long been used for mechanicalto-electrical energy harvesting [5]. However, these materials tend to be stiff and display limitations in mechanical strain abilities. Consequently, they cannot be used for many applications requiring low frequencies and large stroke mechanical excitations (such as human movement). Organic materials, however, are softer and
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more flexible; therefore, the input mechanical energy is considerably higher under the same mechanical force. Piezoelectric polymers, such as PVDF, unfortunately have a much lower piezoelectric coefficient compared to piezoelectric ceramic materials. A study has shown that the energy harvesting is lower than with piezoelectric ceramic bimorphs [6]. Electrostatic-based systems, such as dielectric elastomers, usually require a very high electrical field intensity (50–120 MV/m) to achieve a significant energy harvesting [4]. Recent research has shown that polyurethane, an electrostrictive polymer, is capable of generating strains above 10% under a moderate electrical field (20 MV/m), thus leading to them being considered as potential actuators. Furthermore, these materials are lightweight, very flexible, have low manufacturing costs and are easily molded into any desired shape. Although not very well-known, these materials can also be used for mechanical-to-electrical energy harvesting. Electrostriction is generally defined as a quadratic coupling between strain and electrical field. Assuming a linear relationship between the polarization and the electrical field, the strain Sij and the electric flux density Di are expressed as independent variables of the electrical field intensity Ek, El and stress Tkl by the constitutive relations presented below [7]. ( Sij ¼ Mijkl Ek El þ sEijkl Tkl ð8Þ Di ¼ eTik Ek þ 2 Mijkl El Tkl Here, sEijkl is the elastic compliance, Mijkl is the electrical-field-related electrostriction coefficient, and eTik is the linear dielectric permittivity. There exist two methods for harvesting energy using an electrostrictive material. The first one consists in realizing cycles and the second involves working with the pseudo-piezoelectric behavior.
2.2.1 Energy harvesting cycles The method proposed by Liu et al [8] was inspired by the approach for harvesting energy in the case of dielectric elastomers. The mechanical-to-electrical energy harvesting in electrostrictive materials can be illustrated by for instance the mechanical stress/strain and electric field/flux density plots, such as those shown in Fig. 2. Initially the material presented in Fig. 2 had no applied stress, but then a stress was applied and the state traveled along path A. The applied stress was
Fig. 2 Energy harvesting cycle
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Fig. 3 Energy harvesting cycles under constant electrical field conditions as the material is stressed and unstressed
subsequently reduced. Due to the change in the electrical boundary conditions, the contraction path did not follow path A but path B. Both in the mechanical and electrical planes, the material state traversed a closed loop. In the mechanical plane, the rotation designated that the clear energy flow went from the mechanical to the electrical. The area enclosed in the loop of the mechanical and electrical planes was the same and corresponded to the converted energy density in units of J/m3. Ideally the energy harvesting cycle consists of the largest possible loop, bounded only by the limitations of the material. Liu et al. have analyzed electrical boundary conditions that can be applied for optimizing the energy harvesting. They demonstrated that electrostrictive materials have significant electric energy densities that can be harvested. With the electrical boundary conditions investigated in their work, the best energy harvesting density occurred when the electrical field in the material was increased from zero to its maximum value at a maximum stress, and then returned to zero at a minimum stress (Fig. 3). A constant electrical field E0 existed from state 1 to state 2 as the stress was increased to Tmax. From state 2 to state 3, the electrical field was increased from E0 to E1, then kept constant until the stress was reduced from Tmax to 0 from state 3 to state 4. At zero stress, the electrical field was reduced to E0, returning to state 1. In the dielectricfield plot, the paths 1–4 and 2–3 were not parallel, which was due to the stress dependence of the dielectric constant. Ren et al. [7] investigated a means of using this method for harvesting energy. An experiment was carried out under quasistatic conditions of 1 Hz. The electric parameters were E0 = 46 MV/m and E1 = 67 MV/m, and 22.4 mJ/cm3 of energy was harvested for a transverse strain of 2%. If the piezopolymers or piezoceramics were used with a conventional energy harvesting scheme, the energy harvesting was below 5 mJ/cm3 [9].
2.2.2 Energy Harvesting with Pseudo Piezoelectric Behavior Another way for harvesting energy, using electrostrictive polymers, consists in working with a pseudo piezoelectric behavior. For this, the electrostrictive polymer was subjected to a DC biased electrical field. As the polymer was not piezoelectric, it was necessary to induce polarization with a DC bias in order to obtain the desired pseudo piezoelectric behavior [1, 10, 11].
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Fig. 4 The mechanical configuration of an electrostrictive polymer
An isotropic electrostrictive polymer film contracts along the thickness direction and expands along the film direction when an electrical field is applied across the thickness, assuming that only a nonzero stress is applied along the length of the film (Fig. 4). The constitutive relation, expressed by Eq. 8, can thus be simplified according to ( S1 ¼ M31 E32 þ sE11 T1 ð9Þ D3 ¼ eT33 E3 þ 2 M31 E3 T1 The current induced by the transverse vibration is measured as Z oD3 dA I¼ ot
ð10Þ
A
where A corresponds to the area of the EPC. The current produced by the polymer can thereby be related to the strain and electrical field by # Z " 2 2 M31 oSot1 E3 oE3 T 2 M31 S1 6 M31 E32 e þ I¼ þ dA ð11Þ ot 33 sE11 sE11 A
where oE3=ot and oS1=ot are the time derivates of the electrical field and strain, respectively. Since a DC electrical field (Edc) is applied on the sample, so that oE3 =ot ¼ 0; the short circuit current is therefore given by:
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Z
oS1 dA ot
ð12Þ
A
with 1 sE11 ¼ Y: Here, Y is the Young modulus. Assuming a constant strain, the relation between the displacement and strain S1 in the polymer can be expressed by Eq. 13. S1 ¼
DL u ¼ L0 L0
ð13Þ
Here, DL = u is the amplitude of the transverse displacement, and L0 is the initial value of the length. The electric impedance of a polymer vibrating at a given frequency could be modeled by an equivalent electric circuit. Figure 5 displays the most commonly adopted form of an electrical scheme, where Cp is the capacitance of the clamped polymer and Rp(x) is a resistance representing the dielectric losses [11]. Both of these factors are functions of the frequency caused by the relaxation phenomenon. tan d ¼
1 Rp ðxÞ Cp x
ð14Þ
The first branch from the left part of Fig. 5 is the motional branch, modeled by the current source I0, given in Eq. 12. This equation can be used to model the harvested current from vibrations. A previous study has demonstrated that it was possible to neglect Rp [11]. The dynamic model of the current can thus be simplified to Ih ¼ aðEdc Þ
ou oV þ Cp ot ot
ð15Þ
dc with aðEdc Þ ¼ 2M31LYAE and where u is the displacement. This expression is 0 similar to the typical system of equations in the case of piezoelectricity [12]. This last remark is very interesting. In fact, the nonlinear approach of piezoelectric energy conversion leads to three different high-performance techniques. Badel et al. [13] presented one technique with an approach based on a theoretical model and experimental results. They showed that the harvested power may be
Fig. 5 The equivalent electric circuit of an electrostrictive polymer
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increased by a factor of 8 as compared to the most frequently used techniques. Ongoing work aims at extending this nonlinear approach to electrostrictive materials. Between dielectric elastomers and electrostrictive polymers, the difference in energy harvesting is not very important. The great advantage of a dielectric elastomer is that it is possible of have a strain of more than 300%, to be compared to that of an electrostrictive polymer which is 100%. Nevertheless, these materials require higher electrical fields, causing an increase in the size of the device as well as in the loss due to the conversion of the high voltage.
3 Energy Harvesting Applications using EAPs One of the most important trends in the electronic equipment technology from its origins has been the reduction in size and the increase in functionality. Nowadays, small, handheld, though very powerful devices are commercially available and allow the user to play music, to communicate wirelessly or to compute practically everywhere or, in other words, ubiquitously. In the next years, new products will become available, providing vision and other extended functions to the user. The size of such devices is becoming so small that they are becoming referred to as wearable instead of portable. They can be integrated in objects for everyday use, such as watches, glasses, clothes, etc. EAPs are well suited for harvesting energy from human motion. Natural muscles, the driving force for human motion, are typically of low frequency and intrinsically linear; two characteristics where EAPs offer advantages. Human walking is a good example of an application of energy harvesting using human motion. Proof-of-principle heel-strike generators have been built using dielectric elastomer devices. With more development, it will likely become feasible to obtain up to approximately 1–2 J per step. The upper limit for energy harvesting without causing discomfort to the walker is estimated at 2–4 W for two shoes, assuming a typical 1-Hz step frequency (Fig. 6) [14]. This amount of power is adequate for portable applications such as cell phone battery charging, and emergency locators for soldiers, as illustrated in Fig. 7. The figure also provides the orders of
Fig. 6 A dielectric elastomer heel strike generator [14]
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Fig. 7 Typical power consumptions [15]
magnitude of the powers consumed by various CMOS (complementary metal oxide semiconductor) electronic equipment that could be powered by energy harvesting devices [15]. EAPs have demonstrated excellent performances, and numerous applications appear feasible, but challenges remain. EAPs appear most advantageous for applications requiring low or variable frequencies, low cost, and large areas. Currently, piezoelectric (PZT) materials are the most popular for harvesting mechanical energy because of their compact configuration and compatibility with MEMSs (micro-electro-mechanical systems). Nevertheless, there are inherent limitations such as aging, depolarization, and brittleness. In order to overcome these limitations, the use of EAPs seems promising and exciting.
4 Examples of Material Development 4.1 Introduction The development of one type of electrostrictive polymer with conductive fillers is currently ongoing. The aim is to fabricate an intermediate material between PZT and actual EAPs, i.e., a novel EAP capable of generating strains larger than those of PZT (2%) and, at the same time, requiring low electric driving fields (10–20 MV/m). The present section presents the ongoing investigation as one example of material development within this field.
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4.2 Drawbacks and Solution Two problems were considered: (1) a high driving energy and (2) aggregation. As mentioned above, the strain level of EAPs is quite impressive and it is true that it has significantly increased the potential of these materials. Nevertheless, one should note that these values have been obtained under a high applied electric field, i.e., 120–150 MV/m. Many factors have to be considered when selecting the actuators; e.g., the response speed, the endurance limit or the generated force. It is clear that the advantages and drawbacks can depend on the applications. For the case when the EAPs are intended to be utilized with other electric components, the high electric field requirement becomes a big obstacle. Since PZT needs only a few MV/m at maximum, its replacement by EAPs is limited. Consequently, a reduction of the induced electric field is expected. As an alternative to using the pure polymer material by itself, some micro- or nano-sized fillers can be added to the polymer matrix in order to increase its capacity. Consider the incorporation of ferroelectric fillers, such as PZT particles, into a polymeric matrix. The relative permittivity er of the matrix is small enough to be ignored as compared to that of the ferroelectric fillers. Thus, the strain of such a composite has a direct relation to the permittivity of the fillers. The permittivity of the composite has an exponential relation to the volume of the fillers. In order for the composite to generate a large strain, a high volume percentage of fillers is necessary. Here, a notorious problem appears: aggregation. As the volume percentage of filler particles is raised, aggregation occurs with increasing ease. Also in the case where the diameter of the fillers is lowered, aggregates form more easily due to surface effects becoming increasingly significant. The existence of large aggregates raises the risk of electric break down. Percolation threshold must thus be considered. Numerous researchers have attempted to solve this problem by both chemical and mechanical approaches, such as sol-gel processing, melt-mixing methods, or ultrasonic activation. Although some of these alternatives are quite effective, they often cause the process to become long and complicated. To overcome these problems, a first attempt involved reducing the amount of fillers. This was not possible with ferroelectric fillers but with conductive fillers, and the concept differs from that of ferroelectric fillers. In the case of the conductive fillers, each particle is believed to become a dipole when an electric field is induced on the composite film, and a high volume percentage is essential in order to obtain a reasonable strain. Instead, in the case of conductive filler particles, a small volume percentage is assumed to be sufficient for generating a reasonable strain. It is believed that the electrons remain on the surface of each conductive particle during electric field induction on the composite film since the electric field is null inside. Consequently, a smaller electric field should be enough to drive such composite films with conductive fillers. With a low volume percentage of fillers, the risk of electric short circuiting caused by aggregates should decrease. In the ongoing study, carbon black was chosen as the filler material. Moreover, the micellar form of carbon black particles was selected in order to more efficiently
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avoid the aggregation problem. These micelles were obtained directly in the form of a carbon black ink.
4.3 Experimental Setup The chosen materials were the following: commercially available polyurethane (PU) granules, and a polyether-type thermoplastic TPU5888 from Estane, used as the matrix. There were two types of carbon black (CB), of which the first was CB in micellar form, denoted ‘CB-ink’ here. The diameter of the CB micelles was around from 15 to 30 nm as obtained from transmission electron microscope (TEM) observation [16]. For the sake of comparison, simple nano-sized CB particles, denoted ‘CBP’, were also employed. These CBPs were amorphous carbon nanopowders obtained from Aldrich, with an average particle size of 30 nm. N,N-dimethylformamide (DMF) was used as a solvent for PU. The composite films were prepared by a simple solution casting method [17–19]. PU granules were completely dissolved in DMF at around 75°C. Subsequently, either CB-ink or CBP was added to the solution under non-magnetic stirring. The volume percentage of fillers did not exceed 2% in either of the two cases. At a constant temperature, the stirring time varied between 1 and 2 h. To visualize the dispersion effect, an ultrasonic treatment was carried out on certain solutions. The application time ranged from 2 to 4 h. The mixed solution was then poured onto a glass plate and dried at 60°C under air during 15 h. The films, after being removed from the glass plate, were dried again at 130°C under air. The film thicknesses varied between 7 and 143 lm. Samples without any additives, denoted ‘Pure PU’, were also prepared for comparison. The films were cut into discs with a diameter of 25 mm. For electromechanical characterization, gold electrodes (6 mm in diameter and approximately 20 nm thick) were sputtered onto the two surfaces. The field-induced thickness strain was measured with a double beam laser interferometer (Agilent 10889B) with a precision on the order of 10 nm. Microorder strain could be measured with more precision with two beams as opposed to just a single one. The corresponding bipolar voltage was delivered by an Agilent 33250A function generator, amplified by a factor 1000 through a high-voltage lock-in Trek 10/10B amplifier.
4.4 Results 4.4.1 Dispersion The pure PU sample was completely transparent. If care was not taken, numerous aggregates were easily formed with the CBP addition. This film was basically transparent but many black aggregates could be visualized with the naked eye, as shown in the photographs of Fig. 8a. After ultrasonic activation, large aggregates
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Fig. 8 Photographs of CBink/PU samples: a without aggregates and b with aggregates
were no longer observed. The well-mixed specimen was completely black and the color deepened as the percentage of CB increased. Figure 8b shows a photograph of the CB-ink/PU sample. In this case, even without ultrasonic treatment, no large aggregations could be observed. It was thus possible to readily avoid large aggregates by using CB-ink in micellar form. Figure 9 shows SEM micrographs of the CB-ink/PU samples: (a) and (b) display the samples with 1.7 vol. % of CB-ink, and (c) and (d) present specimens with 0.1 vol. %. The micrographs in Fig. 9a, c demonstrated very good dispersions (some specks of dusts were present on the surfaces), with the exception of a few blobs as the one encircled in Fig. 9a. Figure 9b shows a close-up of this micaceous aggregate. As can be seen in Fig. 9c, the existence of these blobs was very rare in the 0.1-vol. % specimens. Nevertheless, Fig. 9d shows a close-up a one of these rare blobs. These micaceous aggregates were formed with small particle-like spots. Although difficult to properly visualize in the printed photos, the surfaces of these spots did not seem attached to each other. Rather, interfaces were observed. This Fig. 9 SEM micrographs of CB-ink/PU samples: a and b samples with 1.7 vol. % of CB-ink; c and d samples with 0.1 vol. % CB-ink
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signifies that they were some kind of primly particles. The aggregates here were different from those that are often found in polymer composites with particles, as shown in Fig. 8b. In this case, the particles were attached to each other. These micaceous aggregations with CB-ink were denoted ‘‘clusters’’ in order to differentiate between them. It is however unclear whether the electrons remained on each individual surface of the particles in these clusters or on the surface of the whole cluster. Since the electric break down did not occur easily, it was determined that proper electrical contacts could not be made at such low volume percentages.
4.4.2 Overview of Typical Electrical Field-Inducted Measurement Results The results of the pure PU specimens and those with addition of CB showed the same tendency. First, typical results are shown after which the differences will be discussed. Some figures also display simulated data for the sake of comparison. This is discussed in the next section. The induced electrical field corresponded to 2 cycles of a triangle-shaped wave of 0.1 Hz. Each field started from zero and the maximum was varied as 1, 2, 3, 4, 5, 7, 8, 9, 10, 15 and 20 MV/mm. Figure 10 shows the typical result of each strain S behavior versus time. An image of the electrical field is presented for comparison. Most of the time, the more the maximum induced electrical field was raised, the more the maximum strain increased, but not always. The convergence of the maximum strain appeared with certain values of the electrical field. Figure 11 displays the maximum strain as a function of the maximum applied electrical field. The one shows an example of convergence and the other displays non-convergence. The convergence and decrease of the strain can be clearly seen. If the amplitude of the electrical field is increased, it is estimated that also these specimens would demonstrate the convergence. The reason for this will be explained later on. Figure 12 illustrates the typical result of the polarization P as a function of the electric field. The curves are closed and present an almond shape. The appearance of a hysteresis can be seen.
Fig. 10 Typical strain behaviors versus time under 2 cycles of a triangle-shaped electrical field
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Fig. 11 Typical results of the maximum strain versus the maximum applied electrical field
Fig. 12 A typical result of polarization versus the applied electrical field
The typical change of the current and polarization when increasing the maximum electrical field is shown in Fig. 13. The data was obtained at (a) 1 MV/m, (b) at 5 MV/m and (c) at 20 MV/m of the maximum electrical field induction. Figure 14 displays the results of the maximum strain vs. the maximum electrical field for several types of samples: 1.7 and 0.1 vol. % of CB-ink/PU, 2 and 1 vol. % of CBP/PU and 2 samples of pure PU. The conversion of strain for the pure PU appeared to be less than 10 MV/m. Consequently, the strain level was limited to 10%. The samples to which CBP had been added also displayed convergence. These values of maximum strain were slightly higher than 10%. It is clear that the CB-ink/PU samples presented higher strain levels than their pure PU and CBP/PU counterparts. The samples with 1.7 and 0.1 vol. % exhibited maximum strains around 24 and 12%, respectively. It was quite remarkable that even the addition of merely 0.1 vol. % could give rise to higher strain levels than the 2-vol. % CBP/PU sample. Another significant result is that there was a dependence of the maximum strain on the thickness. Figure 15 shows this dependence for the specimen with 1.7 vol. % of CB-ink/PU.
298 Fig. 13 The change of the form of the current and polarization with a maximum electrical field of a 1 MV/m, b 5 MV/m and c 20 MV/m
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Fig. 14 The maximum strain as a function of the maximum applied electrical field Fig. 15 The dependence of the maximum strain on the thickness
4.5 Discussion There are two possible explanations for the convergence of strain: a mechanical and an electrical explanation. The former corresponds to an effect of the measurement. When a triangular electrical field is induced, whether positive or negative, there occurs a deformation of the thickness in the direction of compression. The strain in the present study was based on the compressed deformation. Simultaneously, the elongation occurred in the direction of the diameter of the disk-shaped specimen. This elongation, stimulated by the electric field, was however limited in the elastic domain. It differed from the force application measurement, and it is thus natural that the conversion occurred also with regard to the thickness deformation. Still, a convergence around 10% of strain (pure PU) seemed quite low when caused by such a phenomenon. Another explanation could be an electrical reason, i.e., the change in material characteristics caused by the induced electrical field. The basic constitutive equations are the following:
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S3 ¼ s33 s3 þ a33 P23 E ¼ bP þ 2asP
ð16Þ
Under no-stress condition, taking P ^ eE at a low electrical field regime, the relation between strain and electrical field is always found to be quadratic according to S33 ’ ae2 E2 ;
ð17Þ
where a is an electrostrictive coefficient. Besides, a polymeric composite can be considered as a capacitive material. The current is expressed as _ I ¼ C V;
ð18Þ
where C is the capacitance. With a triangle-shaped electrical induction, the current becomes rectangular as shown in Fig. 13a. The polarization takes on the triangular shape.. Taking into account electrical losses, the current can be written V I ¼ CV_ þ R
ð19Þ
The electrical loss appears R as the inner surface of Fig. 12 since the polarization P is calculated by P ¼ A1 Idt: A represents the conductive surface area and its change is neglected here. Equation 19 does not present the experimental evaluation of Fig. 13. To overcome this problem, polarization was considered. The evolution of polarization is difficult to observe experimentally because of the prevention by the loss, see in Fig. 12. Here, the polarization undergoes saturation according to the following equation: E ð20Þ P ¼ eEsat tanh Esat where Esat is the convergence value of the electrical field. This assumption of polarization was confirmed by the following two conditions: (1) when the electric field is small compared to Esat, P = eE and (2) when it is much bigger, i.e., E Esat, P = eEsat. The strain can be simply expressed by the equation of polarization as S = aP2. Figure 10b is the simulation curves. The form change of strain from a triangular shape at a low electrical field to a round shape at a high electrical field is accurately described with this simulation. Since the loss is not yet concerned, the strain returns to zero at each E = 0 in Fig. 10b. From Q = PA, the current can be also obtained in equation. The simulation resulting current has a nonlinear relation with the electrical field. The simulation curves at both a low and high electrical regime show good correspondence with experimental data.
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5 Example Study of Energy Harvesting using EAPs 5.1 Principle of Measurement of Harvested Power Figure 16 provides a schematic representation of the setup developed for characterizing the power harvested by the polymer film. The mechanical system consisted of a shaker and a capacitive sensor. The shaker produced the vibration force in sinusoidal form, causing the sample to undergo a transverse vibration. The capacitive sensor (Fogale MC 940) measured the transverse displacement of the sample from which the strain S1 was calculated. The electrostrictive polymer was subjected to a DC biased electrical field, produced by a function generator and amplified by the Trek Model 10/10 High-Voltage Power Amplifier. As the polymer was not piezoelectric, it was necessary to induce a polarization with a DC bias in order to obtain a pseudo-piezoelectric behavior. The EPC was excited both electrically and mechanically, in order for its expansion and contraction to induce a current measured by the current amplifier (Keithley 617), thus giving rise to ‘‘an image’’ of the power harvesting by the polymer, due to electrical resistance (Rc). In this setup, the current was chosen as it is known to be less sensitive to the noise from the electrical network (50 Hz) and in order to avoid problems of impedance adaptation. All the data was monitored by an oscilloscope (Agilent DS0 6054A Mega zoom).
Fig. 16 A schematic of the experimental setup for the energy harvesting measurements
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5.2 Materials Two types of commercially available polymers were used: polyurethane and nylon. Moreover, two composites were synthesized specifically for the study. These polyurethane composites were prepared in the laboratory, using a thermoplastic polyurethane, the 58887 TPU elastomer (Estane), as the matrix. Neat polyurethane (PU) films as well as their filled counterparts were prepared by a solution casting method as mentioned in the previous section. Two types of inorganic fillers, a carbon black nanopowder (CB) and silicon carbide (SiC) nanowires, were chosen. Before the mixing, they were individually and ultrasonically dispersed in DMF.
5.3 Results Figure 17 presents the power versus electrical field, for a constant electrical load and strain (S1 = 0.25%), in the case of nylon. The RMS power harvested on the load was derived using P = Rc.I2h. As expected from the model, a quadratic dependence between the power and the static electrical field strain was observed. There was a good agreement between the experimental results and the model, thus validating the developed model for evaluation of the harvested power. Table 1 gives the harvested power (Pharvested_AC = Rc.I2h) measured on the various materials (with a transverse strain of S1 = 0.25% at 5 V/lm and 100 Hz) for a matched load equal to Rc ¼ Cp1x ¼ Zoptimal : In fact, according to [11], there should be an optimal load resistance for which the conversion power is at a maximum. The power density for the polymer and composite was very low. Although this could be considered a disappointing result, it should be kept in mind Fig. 17 Power harvested as a function of the electrical field for a constant strain of 0.25% at 100 Hz, and for an electric load of 2.5 MX
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Table 1 Material characteristics and harvested power density at 100 Hz for a static electrical field of 5 V/lm and a transverse strain of 0.25% Type er Y (MPa) Power harvested (lW/cm3) Neat PU PU 0.5 wt%SiC PU Vol1%CB Nylon
4.8 5.6 8.2 12
40 70 40 2800
0.2 0.4 1.0 2.3
that in the electrostrictive case, the power was proportional to the square of the bias field, evaluated in Table 1 for a relatively low bias value. For example, doubling the value of the bias field to 10 V/lm (which is still quite low) would result in a power that was 4 times larger. The nylon provided the higher harvested power density. According to [11, 20–21], the M31 coefficient was proportional to e0 ðer 1Þ2 =ðY er Þ: In other words, M31.Y was roughly proportional to er and in order to increase the energy harvesting of the polymer, it would be necessary to employ a composite with a high dielectric permittivity. The incorporation of fillers provides extreme interests. As shown in Table 1, the addition of conductive nanofillers rendered it possible to increase the permittivity, which in turn led to an increase in the harvested power.
6 Conclusions This chapter has described the use of EAPs for energy harvesting actuator purposes. After an introduction of some EAPs adapted for energy harvesters, two experimental investigations were presented. Elastrostrictive polymers were mainly introduced with our development and analysis study. One of the problems of EAPs today is their high electric requirement. In order to overcome this problem, the first study was devoted to fabricating an intermediate material between a piezo element and actual EAPs, i.e., an EAP that can exhibit a large strain with a reasonably low electrical field. Carbon black was chosen as a conductive filler. In order to avoid aggregation, a micellar type of carbon black, directly obtained from ink, was used. As a result, more than 20% strain was easily obtained at less than 20 MV/m. The conductive mechanism was investigated. The achievement of large-displacement EAPs requiring a supply energy as low as that of peizo elements should lead the way to a large variety of applications. Furthermore, a second study involving energy harvesting through the use of EAPs was also presented. The obtained results are not yet gratifying, but there is hope that they will be improved soon. EAPs is a very interesting and promising material class, and it is believed that the amount of research devoted to it, and consequently the applications derived from it, will expand from one day to the next.
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References 1. Lebrun, L., Guyomar, D., Guiffard, B., Cottinet, P.-J., Putson, C.: The characterisation of the harvesting capabilities of an electrostrictive polymer composite. Sens. Actuators A Phys. 153, 251–257 (2009) 2. Bar-Cohen, Y.: Electroactive Polymer (EAP) Actuators as Artificial Muscles (reality, potential, and challenges). SPIE Press, Bellingham (2001) 3. Kofod, G., Kornbluh, R., Pelrine, R., et al.: Actuation response of polyacrylate dielectric elastomers. In: Bar-Cohen, Y. (ed.) Smart Structures and Materials: Electroactive Polymer Actuators and Devices. Proc. SPIE, vol. 4329 (2001) 4. Pelrine, R., Kornbluh, R., Eckerle, J., et al.: Dielectric elastomer: generator mode fundamentals and applications. In: Bar-Cohen, Y. (ed.) Smart Structures and Materials: Electroactive Polymer Actuators and Devices. Proc. SPIE, vol. 4329 (2001) 5. Lefeuvre, E., Badel, A., Richard, C., Guyomar, D.: High efficiency piezoelectric vibration energy reclamation. In: Proc. SPIE SSM Conference, San Diego (2005) 6. Liu, Y., Ren, K., Hofmann, H.F., Zhang, Q.M.: Electrostrictive polymer for mechanical energy harvesting. In: Proc. SPIE, Int. Soc. Opt. Eng., vol. 5385, pp. 17–28 (2004) 7. Ren, K., Liu, Y., Hofmann, H., Zhang, Q.M.: An active energy harvesting scheme with an electroactive polymer. Appl. Phys. Lett. 91, 132910 (2007) 8. Liu, Y., Ren, K.L., Hofmann, F., et al.: Investigation of electrostrictive polymers for energy harvesting. IEEE UFFC 52(12), 2411–2417 (2005) 9. Poulin, G., Sarraute, E., Costa, F., et al.: Generation of electrical energy for portable devices: comparative study of an electromagnetic and piezoelectric system. Sens. Actuators A Phys. 116(3), 461–471 (2004) 10. Guyomar, D., Lebrun, L., Putson, C., Cottinet, P.-J., Guiffard, B., Muensit, S.: Electrostrictive energy conversion in polyurethane nanocomposites. J. Appl. Phys. 106, 014910 (2009) 11. Cottinet, P.-J., Guyomar, D., Guiffard, B., Putson, C., Lebrun, L.: Modeling and experimentation on an electrostrictive polymer composite for energy harvesting. IEEE Trans. Ultrason. Ferroelectr. Freq. Control 57(4), 774–784 (2010) 12. Badel, A., Guyomar, D., Lefeuvre, E., et al.: Efficiency enhancement of a piezoelectric energy harvesting device in pulsed operation by synchronous charge inversion. J. Intell. Mater. Syst. Struct. 16, 889 (2005) 13. Badel, A., Benayad, A., Lefeuvre, E., Lebrun, L., Richard, C., Guyomar, D.: Single crystals and nonlinear process for outstanding vibration-powered electrical generators. IEEE Trans. Ultrason. Ferroelectr. Freq. Control 53, 673–684 (2006) 14. Prahald, H., Kornbluh, R., Pelrine, R., et al.: Dielectric elastomers and their applications in distributed actuation and power generator. In: Proceedings of ISSS 2005. International conference on Smart Materials Structures and Systems, India (2005) 15. Sebald, G., Lefeuvre, E., Guyomar, D.: Pyroelectric energy conversion: optimization principles. IEEE Trans. Ultrason. Ferroelectr. Freq. Control 55(3), 538–551 (2008) 16. Kanda, M., Yuse, K., Guiffard, B., Guyomar, D.: Large field induced strain in carbon nanofilled composite polyurethane (PU). In: Proc. Smart’09, 493 pp (2009) 17. Guiffard, B., Seveyrat, L., Sebaid, G., Guyomar, D.: Enhanced electric field-induced strain in non-percolative carbon nanopowder/polyurethane composites. J. Phys. D 39, 3053–3057 (2006) 18. Petit, L., Guiffard, B., Seveyrat, L., et al.: Actuating abilities of electroactive carbon nanopowder/polyurethane composite films. Sens. Actuators A Phys. 148, 105–110 (2008) 19. Guyomar, D., Matei, D.F., Guiffard, B., Le, Q., Belouadah, R.: Magnetoelectricity in polyurethane films loaded with different magnetic particles. Mater. Lett. 63, 611–613 (2009)
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20. Lebrun, L., Guyomar, D., Guiffard, B., Cottinet, P.-J., Putson, C.: The characterization of the harvesting capabilities of an electrostrictive polymer composite. Sens. Actuators A Phys. 153, 251–257 (2009) 21. Guyomar, D., Lebrun, L., Putson, C., Guiffard, B., Cottinet, P.-J., Muensit, S.: Electrostrictive energy conversion in polyurethane nanocomposites. J. Appl. Phys. (2009). doi:10.1063/1.3159900
Elastomeric Nanocomposites for Aerospace Applications James Njuguna, Krzysztof Pielichowski and Agnieszka Leszczyn´ska
Abstract Nanotechnology adds a new capability to the design of materials resulting in significant performance improvements at macroscopic level. In particular, nanostructured materials open a new paradigm wherein fibres and matrix resins can be tailored to enhance composite properties of interest; just as fibres and ply orientation is currently used to tailor current advanced composites without increasing the mass. This chapter focuses on these new materials with improved electrical, thermal, and mechanical properties. Special attention is given on elastomeric nanocomposites from polyamides, polyurethanes, polyaniline and polyethylene terephthalate due to their diversified applications. Polyimide, polyarylacetylene poly(aryl–ether–ether–ketone) and poly(p-phenylene benzbisoxazole) nanocomposites are also covered for high performance and temperature aerospace applications. Exploitation of these new elastomers’ intrinsic properties relies on the ability to systematically synthesize, characterize, and integrate standardized materials into actual commercial products.
1 Introduction The uses of polymer nanocomposites in structures have several predictable impacts on aerospace design and applications, primarily by providing a safer,
J. Njuguna (&) Department of Sustainable Systems, Cranfield University, Bedfordshire MK43 0AL, UK e-mail:
[email protected] K. Pielichowski and A. Leszczyn´ska Department of Chemistry and Technology of Polymers, Cracow University of Technology, ul. Warszawska 24, 31-155 Kraków, Poland
V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_12, Ó Springer-Verlag Berlin Heidelberg 2011
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faster, and eventually cheaper transportation in the future. The most obvious of which is significant airframe weight reduction stemming from polymer nanocomposites low density and complemented by its high strength and modulus. Proper engagement of the nanoparticles in composite systems depends strongly on the ability to homogeneously disperse them throughout the matrix without destroying their integrity. The properties of the matrix, the distribution and properties of the filler as well as the nature of their interface control the behaviour of a typical composite material [1–3]. Thus, the nanoparticles often strongly influence the properties of the composites at very low volume fractions. This is mainly due to their small interparticle distances and the transformation of a large fraction of the polymer matrix near their surfaces into an interphase of different properties as well as to the consequent change in morphology [4]. As a result, the desired properties are usually reached at low filler volume fraction, which allows the nanocomposites to retain the macroscopic homogeneity and low density of the polymer. Besides, the geometrical shape of the particles plays an important role in determining the properties of the composites. Consequently, nanocomposites have attracted much scientific and industrial interest in recent times. So far, most of the scientific work has been focussing on the synthesis of polymer nanocomposites and on the study of their physical and mechanical properties. The use of these nanocomposites as matrix in fibre-reinforced composites is in its infancy. Nano-composites from polymeric matrix materials (thermoplasts or thermosets) reinforced with nano-sized fillers (see Fig. 1) such as carbon nano-tubes, nano-sized metal and metal oxide particles or lamellar inorganic or carbonaceous nanoparticles are an active area of research [5, 6]. The primary aim is to improve the deficiencies of existing composities or to extend their functional limits. For example, the improved mechanical performance of polymeric nanocomposites is macroscopically depicted in the increase in for instance, fracture energy values. Nanofillers, due to their size, can be significantly present in the plastic zone deformation while in the case of micro-particles only few of them participate in this small zone. In this way, nanofillers can lead to increased fracture properties of the brittle matrix. On the other hand, the matrix material is pointed out to be a key parameter for [2, 3] mode I delamination resistance (critical strain energy release rate for delamination in an opening mode) of fibre reinforced polymers. Therefore, enhancement in the matrix fracture toughness can lead to an overall advanced fracture behaviour. Additionally, the influence of the nanofibres has also been extended to the reinforcing fibres by making more nanofibres to be involved during the delamination process and thus increasing the fracture toughness. Further, by invoking the properties of nanostructures it may be possible to control the shock wave propagation in the material and thus, significantly enhance the impact energy dissipation; however, only limited research works to date concern about damping composite structures and shock resistant nanomaterials. Adsorption of vibration energy by mechanical damping is significant problem in many engineering designs; high damping materials are used to reduce vibration in aircraft and other machinery. The benefits of damping treatment are advanced durability, reliability and service life of
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Fig. 1 a Comparison of conventional fiber composite and nanocomposites. b Logarithmic isolines of interfacial (surface) area/volume of particles (lm-1 = m2/ml) with respect to the aspect ratio, a = H/R, and largest dimension of particle (R = radius, H = height, length) based on approximating particles as cylinders (area/volume = 1/H ? 1/R). Aspect ratios greater than one correspond to rods (length/diameter) and less than one to plates (height/diameter) (reproduced from Refs. [5, 6])
components, reductions in weight, noise and costs. It should be noted, however, that damping and integrity often have opposite requirements (see Fig. 2) [7]; viscoelastic polymers are good for damping, but stiffness is decreased at higher temperatures.
2 Polymer Nanocomposites A key advantage of the use of nanocomposite instead of other microparticulate fillers to improve the fibre composite properties is that the properties can be improved without any change in the processing conditions. The most promising current approaches towards increasing the orientation of nanoscale reinforcements within a matrix include optimisation of the extrusion die and stretching the composite melt to form films and fibres. One complication is that the microstructure of semicrystalline polymer matrices is influenced not only by the processing history but also by the presence of nanoparticles. The addition of various types of nanofillers to polymers has already been observed to influence the
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Fig. 2 Stiffness-damping map for nanoparticle-reinforced and conventional composites (reproduced from Ref. [7])
crystallisation kinetics and resulting morphology. Such changes in matrix morphology need to be considered when evaluating the nanocomposite performance with regard to the intrinsic filler properties. The effects of carbon nanotubes or nanofibres on such oriented polymer systems, although significant, have not yet been fully established. Finally, it should be noted that the presence of additives such as colouring pigments has been shown to influence matrix morphology during fibre spinning, whilst there is the whole technology of nucleating agents which are deliberately added to influence crystalline microstructure. Nanoparticle reinforcement of fibre composites has been shown to be a possibility, but much work remains to be performed in order to understand how nanoreinforcement results in major changes in material properties. The understanding of these phenomena will facilitate their extension to the reinforcement of more complicated anisotropic structures and advanced polymeric composite systems. The property and performance enhancements made possible by nanoparticle reinforcement may be of great utility for fibre-reinforced composites that are applied in aerospace industry.
2.1 Polyamide (PA) Nanocomposites In the early 1990s of twentieth century, Toyota Research group synthesized PA-6based clay nanocomposites that demonstrated the first use of nanoclays as reinforcement of polymer systems. They concluded that nanoclays not only influenced
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the crystallization process, but that they were also responsible for morphological changes. Recognizing these benefits, many researchers, using a variety of clays and polymeric matrices, have produced nanocomposites with improved properties. For instance, Liu et al. [8] reported that there was an increase in storage elastic modulus of 100% when clay content was up to 8 wt% in comparison with net PA11. Usuki et al. [9] polymerized e-caprolactam in the interlayer of an organoclay to form a nanocomposite. This material contained only 4.2 wt% clay and had a 50% increase in strength, an increase in the heat distortion temperature of 80°C, a 100% increase in tensile modulus, and a 20% increase in impact resistance. The electrospun nanocomposite fibres have great potential for the applications where both high surface-to-volume ratio and strong mechanical properties are required such as the high-performance filters and fibre reinforcement materials. Since the mechanical properties of fibres in general improve substantially with decreasing fibre diameter, there is considerable interest in the development of continuous electrospun polymer nanofibres. In this respect, Lincoln et al. [10] reported that the degree of crystallinity of PA-6 annealed at 205°C increased substantially with the addition of montmorillonite (MMT). This implied that the silicate layers could act as nucleating agents and/or growth accelerators. In contrast, the study of Fong et al. [11] showed a very similar overall degree of crystallinity for electrospun PA-6 and PA-6/Cloisite-30B nanocomposite fibres containing 7.5 wt% of organically-modified MMT (OMMT) layers. Fornes and Paul [12] have found that OMMT layers could serve as nucleating agents at 3% concentration in PA-6/OMMT nanocomposite but retarded the crystallization of PA-6 at a higher concentration of around 7%. In addition, the differences in the molecular weight of PA-6 and the solvent used for electrospinning were also expected to have different impacts on the mobility of PA-6 macro chains and the interactions between them and OMMT layers, which may also affect the crystallization behaviour of polymer during the electrospinning. Li et al. [13] manufactured PA-6 fibres and nanocomposite fibres with average diameters around 100 nm by electrospinning using 88% aqueous formic acid as the solvent. The addition of OMMT layers in the PA-6 solution increased the solution viscosity significantly and changed the resulting fibre morphology and sizes. Transmission electron microscopy (TEM) images of the nanocomposite fibres, including ultra-thin fibre sections, and wide angle X-ray diffraction (WAXD) results showed that OMMT layers were well exfoliated inside the nanocomposite fibres and oriented along the fibre axial direction. The degree of crystallinity and crystallite size were both increased for the nanocomposite fibres and more significant for the fibres electrospun from 15% nanocomposite solution, which exhibited the finest average fibres size. As a result, the tensile properties of electrospun nanocomposites were greatly improved. The Young’s modulus and ultimate strength of electrospun nanocomposite fibrous mats were improved up to 70 and 30%, respectively, when compared with PA-6 electrospun mats. However, the ultimate strength of the nanocomposite fibrous mats electrospun from 20% nanocomposites solution was decreased by about 20% due to their larger fibre sizes. The Young’s modulus of PA-6 electrospun single fibres with a diameter around 80 nm was almost double
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the highest value that had been reported for the conventional PA-6 fibres and could be improved by about 100% for the electrospun nanocomposite single fibres of similar diameters. In another interesting study [14], a range of polymer matrices were examined including poly(vinyl alcohol), poly(9-vinyl carbazole) and polyamide. To compare production methods, polymer composite films and fibres were produced. It was found that by adding various mass fractions of nanofillers, both the Young’s modulus and hardness increased significantly for both films and fibres. In addition, the thermal behaviour was seen to be strongly dependent on the nanofillers added to the polymer matrices. Wu et al. [15] prepared carbon fibre and glass fibre reinforced PA-6 and PA-6/clay nanocomposites. The fabrication method involved first mechanically mixing PA-6 and PA-6/clay with E-glass short fibre (6-mm long) and carbon fibre (6-mm long), separately. A twin-screw extruder at a rotational speed 20 rpm extruded the fibres. The temperature profiles of the barrel were 190–210–230–220°C from the hopper to the die. The extrudate was pelletized, dried, and injection moulded into standard test samples for mechanical properties test. The injection-moulding temperature and pressure were 230°C and 13.5 MPa, respectively. The research found out that the tensile strength of PA-6/ clay containing 30 wt% glass fibres was 11% higher than that of PA-6 containing 30 wt% glass fibre, while the tensile modulus of nanocomposite increased by 42%. Flexural strength and flexural modulus of PA-6/clay were found similar to PA-6 reinforced with 20 wt% glass fibres. It was eluded that the effect of nanoscale clay on toughness was more significant than that of the fibre. Heat distortion temperatures of PA-6/clay and PA-6 were 112 and 62°C, respectively. Consequently, the heat distortion temperature of fibre reinforced PA-6/clay system was almost 20°C higher than that of fibre reinforced PA-6 system. Notched Izod impact strength of the composites decreased with the addition of the fibre. The scanning electron microscopy (SEM) microphotographs showed that the wet-out of glass fibre was better than carbon fibre. The study concluded that the mechanical and thermal properties of the PA 6/clay nanocomposites were superior to those of PA-6 composite in terms of the heat distortion temperature, tensile and flexural strength and modulus without sacrificing their impact strength. This was attributed to the nanoscale effects, and the strong interaction force existing between the PA-6 matrix and the clay interface. In case of short fibres, Akkapeddi [16] prepared PA-6 based nanocomposites using chopped glass fibres. In a typical experiment, a commercial grade PA-6 of MW = 30,000 and specially designed functional organo-quaternary ammoniumclay complexes based on MMT or hectorite type clays. Freshly dried PA-6 (moisture \ 0.05%) was blended with 3–5 wt% of a selected organoclay powder and extruded at 260°C in a single step, under high shear mixing conditions. Alternatively, the organoclay was master-batched first into PA-6 (at 25 wt% loading and then re-extruded in a second step with more PA-6 to dilute the clay content to B5 wt%. Conventional chopped glass fibre with 10 lm diameter and about 3 mm length was then added, as an optional reinforcement through a downstream feed port of the twin screw extruder. The glass fibre was compounded
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Fig. 3 Flexural modulus against glass fibre density (GF), PA-6/nanoclay (PA-6/ NC) versus PA-6 moulding resins (reproduced from Ref. [9])
with the molten, premixed PA-6 nanocomposite either as a one-step extrusion process or in a second extrusion step. The extrudate was quenched in a water bath and pelletized. The pellets were dried under vacuum at 85°C, and injection moulded into standard ASTM test specimens. As shown in Fig. 3, significant improvements in modulus were achievable in both the dry and the moisture conditioned state for PA-6 nanocomposites compared to standard PA-6, at any given level of glass fibre reinforcement. In particular, a small amount (3–4 wt%) of nanometer scale dispersed layered silicate was capable of replacing up to 40 wt% of a standard mineral filler or 10–15 wt% of glass fibre to give equivalent stiffness at a lower density. In addition, improved moisture resistance, permeation barrier and fast crystallization/ mould cycle time contribute to the usefulness of such composites. Vlasveld et al. [17, 18] investigated fibre–matrix adhesion in glass-fibre reinforced PA-6 silicate nanocomposites. The main reinforcing phase consisted of continuous E-glass fibres, whereas the PA-6 based matrix was a nanocomposite reinforced with platelets of exfoliated layered silicate. Two different types of nanocomposite were used with different degrees of exfoliation of the silicate layers: one with non-modified silicate and one with an organically modified silicate. They developed nanocomposite laminates by sol–gel and modified diaphragm methods. The route for the preparation of PA-6 nanocomposites consisted of melt-compounding polymer (AkulonÒ K122D) with organically modified fluorine-containing hectorite (SomasifÒ MEE and SomasifÒ ME-100) by means of a co-rotating twin-screw extruder at 240°C. For the SomasifÒ MEE nanocomposite materials, first an 11 wt% MEE master batch was compounded. To obtain the various concentrations of the MEE nanocomposite, the master batch was extruded for a second time without dilution for the 11 wt% nanocomposite, or diluted with PA-6 to concentrations of 6.1 and 2.7 wt%. The 2.5 wt% SomasifÒ ME-100
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nanocomposite material was produced by diluting a 10% organoclay master batch with polymer in the extruder. All mentioned percentages are weight percentages silicate as measured with a thermogravimetric analyser after heating for 40 min at 800°C in air. Two demands for the preparation of the single fibre fragmentation specimen had to be met: the fibre had to lie straight in the centre of the specimen, and the matrix material of the specimen had to be thin enough to be transparent, since the fibre fragments were examined and measured using an optical microscope. A Fontijne hot plate press heated to 240°C was used to produce the films necessary for the single fibre fragmentation test specimen preparation. Single fibres were carefully extracted from a fibre bundle and placed with a distance of approximately 2 cm parallel to each other between the PA or nanocomposite films. The hot plate press at the same temperature was used to melt the polymer films and a pressure of 0.8 N/mm2 was applied for 30 s to provide the necessary bonding with the fibre. After cooling between cold metal plates, tensile test specimens were prepared. It was observed that the ultimate strength and stiffness increased by adding 1% SiO2 nanoparticles, while little improvement in fatigue behaviour was found. It was concluded that the failure mechanism was by interfacial de-bonding and that both the addition of nanoparticles and moisture conditioning had a negative effect on the bonding between the matrix and the glass fibres. In addition, the researchers noted that in the formed composites the adhesion between the nanocomposites and the carbon fibres (Fig. 4) was probably worse than between the unfilled PA-6, reducing the potentially positive influence of the increased matrix modulus. In a parallel research, Vlasveld et al. [18] developed three-phase thermoplastic composite, consisting of a PA-6 nanocomposite matrix and a main reinforcing phase of woven glass or carbon fibres. The nanocomposite used in this research had moduli that were much higher than those of unfilled PA-6, also above Tg and in moisture conditioned samples. Flexural tests on commercial PA-6 fibre
Fig. 4 Flexural strength of carbon fibre composites with PA6, a commercially available PA6 nanocomposite (Unitika M1030D from Unitika) and nanocomposite matrix as a function of the matrix modulus (dry and moisture conditioned) (reproduced from Ref. [18])
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composites showed decrease of the flexural strength upon increasing temperature. The researchers claimed that the strength of glass fibre composite can be increased by more than 40% at elevated temperatures and the temperature range at which a certain minimum strength is present can be increased by 40–50°C. Carbon fibre composites also showed significant improvements at elevated temperatures, although not at room temperature. Based on flexural tests on PA-6 based glass and carbon fibre composites over a large temperature range up to near the melting point, it became clear that for these fibre composites it is important to have a reasonably high matrix modulus: both glass and carbon composites were very sensitive to a decrease of the matrix modulus below values around 1 GPa. At higher moduli, carbon fibre composites are more sensitive to the matrix modulus than glass fibre composites. The modulus of unfilled PA-6 decreased below the (arbitrary) 1 GPa level just above Tg; it is noteworthy that the nanocomposites used in this research had moduli that were much higher and stayed above the 1 GPa level up to 160°C, which was more than 80°C higher than for unfilled PA-6. The nanocomposites also showed much higher moduli in moisture conditioned samples, and even in moisture conditioned samples tested at 80°C the modulus was much higher than that of the dry unfilled PA-6, again well above 1 GPa. Dynamic mechanical analysis (DMA) measurements showed that the nanocomposites did not show a change of Tg in the dry state, and that the reduction of the modulus upon absorption of moisture was due to the Tg decrease. An assessment of reactively processed anionic polyamide-6 (APA-6) for use as matrix material in fibre composites was conducted by van Rijswijk et al. [19], who also compared it with melt processed PA-6 and PA-6 nanocomposites. A special designed lab-scale mixing unit was used to prepare two liquid material formulations at 110°C under a nitrogen atmosphere: a monomer/activator-mixture in tank A and a monomer/initiator-mixture in tank B, as shown on Fig. 5. After individually degassing both tanks (15 min at 100 mbar), the two material feeds were mixed by using a heated (110°C) static mixer and dispensed (1:1 ratio) into a heated (110°C) buffer vessel with nitrogen protective environment. Stainless steel infusion mould (Fig. 3) was used together with a 3 mm thick stainless steel cover plate (not shown) to manufacture neat APA-6 panels (250 9 250 9 2 mm). Homogeneous heating of the mould was obtained by placing it in a vertically positioned hot flat plate press. A silicon tube connected the resin inlet of the mould with the buffer vessel and the resin outlet with a vacuum pump. Infusion from bottom to top was necessary to prevent entrapment of air. A pressure control system was used to precisely set the infusion and curing pressure (absolute pressure in the mould cavity). Loss of control over the pressure in the mould cavity due to solidification of resin in the unheated outlet tube had to be prevented. To avoid this, a buffer cavity was machined in the mould near the outlet to slow down the infusion, hence giving ample time to stop the resin flow before it was able to exit the mould. For every infusion pressure, the infusion time to reach the buffer cavity was determined visually by replacing the steel cover plate by a glass one. Additionally, a resin trap and a cold trap were placed directly after the mould to protect the vacuum pump. The mechanical properties of APA-6 (commercial grade of
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Fig. 5 Infusion equipment from left to right: Mini Mixing Unit ‘‘MMU-TU Delft’’, resin reservoir, stainless steel infusion mould, resin trap, cold trap and vacuum pump (reproduced from Ref. [12])
PA6, AkulonÒ K222D) and HPA-6 (low MW injection-moulding grade hydrolytically polymerised PA-6) nanocomposites were compared with injection moulded neat HPA-6. As expected, the HPA-6 nanocomposite had the highest modulus over the entire range of temperatures (20–160°C) and moisture contents (0–10 wt%) tested. However, APA-6 came close and had the highest maximum strength due to its characteristic crystal morphology, which was directly linked to the reactive type of processing used. The same morphology, it was claimed, also made APA-6 slightly less ductile compared to melt processed HPA-6. Compared to the melt processed HPA-6, APA-6 polymerised at 150°C and the HPA-6 nanocomposite offered a higher modulus at similar temperature, or similar modulus at a higher temperature (40–80°C increase). It is noteworthy that such an increase in maximum use temperature, related to the heat distortion temperature, can seriously expand the application field of PA-6 and PA-6 composites. For all PAs, temperature and moisture absorption reduced the modulus and the strength and increased the maximum strain, which was directly related to the glass transition temperature. Whereas with increasing testing temperature at a certain moment the Tg of the dry polymer was exceeded, moisture absorption reduced the Tg at a certain point below the testing temperature. However, the effect of both was in essence the same. Retention of mechanical properties of APA-6 after conditioning at 70°C for 500 h and subsequent drying was demonstrated. Conditioning submersed in water at the same temperature, however, resulted in a brittle material with surface cracks, as is common to most polyamides due to continued crystallization and removal of unreacted monomer. Given the fact that submersion at elevated temperatures is usually not an environment in which PA-6 and its composites are applied, the encountered property reduction was therefore not
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detrimental for application of these materials. The overall conclusion of the comparative study for application of the polyamides as matrix material in fibre composites was that both APA-6 and the HPA-6 nanocomposites outperformed the melt processed HPA-6 in terms of modulus and maximum strength. Therefore, the researchers concluded that both special grade PAs may be expected to enhance the matrix dominated composite properties like compressive and flexural strength, provided that a strong fibre-to-matrix interphase is obtained. The presence of multi wall carbon nanotubes (MWNTs) improves the thermal stability of PA-6 under air obviously, but has little effect on the thermal degradation behaviour of PA-6 under nitrogen atmosphere. The thermal degradation mechanism of PA-6 has been proposed by Levchik et al. [20]. A comparative study was conducted by Sandler et al. [21] on melt spun PA-12 fibres reinforced with carbon nanotubes and nanofibres. A range of multi-walled carbon nanotubes (MWNT) and carbon nanofibres (CNF) were mixed with a PA-12 matrix using a twin-screw microextruder, and the resulting blends spun to produce a series of reinforced polymer fibres. The work aimed to compare the dispersion and resulting mechanical properties achieved for nanotubes produced by the electric arc and a variety of chemical vapour deposition techniques. A high quality of dispersion was achieved for all the catalytically-grown materials and the greatest improvements in stiffness were observed using aligned, substrate-grown, carbon nanotubes. The use of entangled MWNT led to the most pronounced increase in yield stress, most likely as result of increased constraint of the polymer matrix due to their relatively high surface area. The degrees of polymer and nanofiller alignment and the morphology of the polymer matrix were assessed using X-ray diffraction (XRD) and differential scanning calorimetry (DSC). The carbon nanotubes were found to act as nucleation sites under slow cooling conditions, the effect scaling with effective surface area. Nevertheless, no significant variations in polymer morphology as a function of nanoscale filler type and loading fraction were observed under the melt spinning conditions applied. A simple rule-of-mixture evaluation of the nanocomposite stiffness revealed a higher effective modulus for the MWNT compared to the CNF, as a result of fewer imperfections in graphitic crystall lattice of MWNT. In addition, this approach allowed a general comparison of the effective nanotube modulus with those of nanoclays as well as common short glass and carbon fibre fillers in melt-blended PA composites. The experimental results further highlighted the fact that the intrinsic crystalline qualities, as well as the straightness of the embedded nanotubes, were significant factors influencing the reinforcement capability. In an intumescent ethylene-co-vinyl acetate (EVA)-based formulation, using PA-6 clay nanocomposite instead of pure PA-6 (carbonization agent) has been shown to improve the fire properties of the intumescent blend. Using clay as ‘‘classical’’ filler enabled the same level of flame retardation performance to be obtained in the first step of the combustion as when directly using exfoliated clay in PA6. But in the second half of the combustion, the clay destabilizes the system and increases the flammability. Moreover, a kinetic modeling of the degradation of the EVA-based formulations shows that adding clay to the blend enables same
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mode of degradation and the same invariant parameters as for the polyamide-6 clay nanocomposite containing intumescent blend. The increase in the flammability in the second half of the combustion shows the advantages of using nanoclay rather than micronclay in an intumescent system [20]. Organically modified clay-reinforced polyamide-6 was subjected to accelerated heat aging to estimate its long-term thermo-oxidative stability and useful lifetime compared to the virgin material [22]. Changes in molecular weight and thermal and mechanical properties were monitored and connected to the polymer modification encountered during aging. Generally, the strong interaction between the matrix and the clay filler renders the polymer chains, mainly that adjacent to silicates that are highly restrained mechanically, enabling a significant portion of an applied force to be transferred to the higher modulus silicates. This mechanism explains the enhancement of tensile modulus that the non-aged clay-reinforced PA6 exhibited (1,320 MPa) with regard to the neat polymer (1,190 MPa), as shown in Fig. 6.
2.2 Polyurethane (PU) Nanocomposites The physical properties of PU are derived from their molecular structure caused by interactions between the polymer chains. The segmental flexibility, the chain entanglement, and the cross-linking effects are all factors that influence on the properties and determine the use of the end-products [23]. Thermoplastic PU elastomers (TPUs) are linear block copolymers consisting of alternating hard and soft segments. The hard segments are composed of alternating diisocyanate and chain extender molecules (i.e. diol or diamine), while the soft segments are formed from a linear, long-chain diol. Phase separation occurs in TPUs because of the thermodynamic incompatibility of the hard and soft segments. The segments aggregate into microdomains resulting in a structure consisting of glassy (hard domains) and rubbery (soft domains) states that are below and above their glass transition temperatures at room temperature, respectively. The hard domains gain their rigidity through physical crosslinking (hydrogen bonding between hard segments), and provide filler-like reinforcement to the soft segments [24, 25]. To improve their mechanical properties PU/layered silicate nanocomposites have received increasing interest. For instance, Ma et al. [26] studies on elastomeric PU/clay nanocomposite based on poly(propylene glycol), glycerol propoxylate and toluene diisocyanate (TDI) showed that the gallery distance of the clay in the hybrid was enlarged from 1.9 to 4.5 nm or more. Introducing clay in the PU matrix resulted in an increase in both the tensile strength and elongation at break. When the clay content reached about 8%, the tensile strength and elongation at break were two times and five times higher to that of the pure PU, respectively. In Han et al. [27] work, the tensile properties of the PU/MMT nanocomposites prepared by solution intercalation displayed higher enhancement relative to PU matrix. Xu et al. [28] proved that the silicate layer spacing in the nanocomposites
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Fig. 6 Tensile properties of PA6 and PA6 nanocomposite, oven-aged at a 120°C and b 150°C (reproduced from Ref. [28])
increased significantly compared with the neat organically layered silicate (OLS), signifying the formation of intercalated PU-urea/OLS structures. The nanocomposite materials exhibited increased modulus with increasing OLS content, while maintaining polymer strength and ductility. Zilg et al. [29] prepared PU
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nanocomposites from modified reactive fluoromica clay. The dried organophilic mica was dispersed by means of a high shear mixer in trihydroxy-terminated oligopropyleneoxide (average MW = 3,800 g/mol). Stable and transparent polyol dispersions were obtained and then cured with 4,40 -diisocyanatophenylmethane (MDI) and accelerated with 0.6 wt% N,N-dimethyl-benzyl-amine. The final nanocomposite PUs showed an increase in tensile strength and elongation in conjunction with a slight decrease in Young’s modulus, whereas non-modified mica gave increased stiffness and only a small increase in tensile strength. Segmented PU/clay nanocomposites have also been developed by Chen et al. [30, 31]. The first report by this group described the initial synthesis of a polycaprolactone/ clay (PCL/clay) nanocomposite by in situ polymerisation and its subsequent incorporation into the solution synthesis of a segmented PU as a partial replacement of the chain extender (1,4-butanediol) in the hard block [30]. Later on, the same group developed PU/clay nanocomposites by directly mixing organoclay with PU using the technique of intercalation from dimethylacetamide solution [32]. In both cases the PU/clay nanocomposites developed by Chen et al. showed enhancements in properties such as modulus and strength. Li [33] proved that a complete and effective entry of monomers into the OMMT layers is very difficult when OMMT content is higher. However, Song et al. [34] reported that the tensile strength of PU/OMMT nanocomposites with 4.7 wt% OMMT content was double compared to that of bulk PU, while Tortora et al. [35] used X-ray analysis to show that exfoliation occurred for low MMT content, whereas for higher contents the intercalated clay rearranged to a minor extent. However, either exfoliated or delaminated samples showed an improvement in the elastic modulus and yield stress but a decrease in the stress and strain at breaking on increasing the clay content. Song et al. [36] synthesised high performance nanocomposites comprising a PU elastomer, based on poly(propylene glycol), 4,40 -methylene bis(cyclohexyl isocyanate) and 1,4-butandiol, and an organically modified layered silicate. The tensile strength and strain at break for these PU elastomer nanocomposites increased more than 150%, but the hardness remained unchanged. The fatigue properties were significantly improved too, particularly for the nanocomposites with 3 wt% organoclay. The effects of the isocyanate index on the mechanical properties of the PU elastomer nanocomposites were subjected to investigations and found that an isoyanate index of 1.1 resulted in the best improvement in stress and elongation at break. In another development, nanocomposites with different concentrations of soft segments (50 and 70%) and prepared by adding nanosilica (up to 30%) to the single-phase PU matrix displayed higher strength and elongation at break but lower density, modulus and hardness than the corresponding micron size silicafilled PUs [37]. Corresponding results have been presented elsewhere, whereby the addition of a small amount of nanosilica increased the hardness, abrasion resistance and tensile properties of the polymer films, but similarly the properties worsened at higher nanosilica contents [38]. In a parallel work, a two fold increase in the tensile strength has been reported for 1% PU/benzidine-MMT as compared
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to that of pure PU [31]. Mishra et al. [39] DMA results indicated a significant increase in storage modulus and tensile strength with increased organoclay loading in millable PU/organoclay nanocomposites. In the study, clay treated with methyl tallow bis(2-hydroxyethyl) quaternary ammonium chloride was used as an organoclay for nanocomposite preparation. The mechanical properties improvements were in good agreement with those reported by Kim et al. [40] on nanocomposites synthesized using amphiphilic urethane precursor (APU)/Na+-MMT emulsions with microphase-separated structure that had greater tensile strength than those prepared with melt-mixed APU/Na+-MMT mixtures. The modulus, tensile strength and hardness were enhanced by the reinforcing effect of intercalated or exfoliated organoclay. It was claimed that the reinforcing effect was more evident when the elasticity of the PU matrix was increased at the temperature range above the Tg of hard segments and in the test where the deformation was large [41]. In Wang et al. [42], Fourier transform infrared spectroscopy (FTIR), WAXD and TEM results verified the incorporation of the OMMT (the MMT modified with was N-octadecyl N,N,N-trimethyl ammonium chloride (DK1) and N-12-dihydroksyethyldodecyl-N,N,N-trimethylammonium chloride (DK2)) into EPU matrix and revealed that the degree of basal-spacing expansion was largely increased. The enhanced mechanical and physical properties demonstrated efficient reinforcing and better thermal stability properties of the latter OMMT. The tensile strength of 3% EPU/DK2 was about 350% higher than that of pure EPU and the elongation at break also showed a remarkable enhancement compared with that of pure EPU. When the amount of MMT increased higher than 3%, the tensile strength decreased, which was blamed to the aggregates of modified MMT in the composites as reported elsewhere by Younghoon and James [43]. Nevertheless, this is in contrast to the work results of Ni et al. [44] where OMMT content [3% reportedly resulted to increase in tensile strength and elongation at break. The work synthesized a novel polyether PU/clay nanocomposite was using poly(tetramethylene glycol), 4,40 -diphenylmethane diisocyanate (MDI), 1,6-hexamethylenediamine, and modified Na+-MMT [44]. Here, OMMT was formed by applying 1,6-hexamethylenediamine as a swelling agent to treat the Na+-MMT. The X-ray analysis showed that exfoliation occurred for the higher OMMT content (40 wt%) in the polymer matrix. The mechanical analysis indicated that, when the OMMT was used as a chain extender to replace a part of the 1,2-diaminopropane to form PU/clay nanocomposites, the strength and strain at break of the polymer was enhanced when increasing the content of OMMT in the matrix. When the OMMT content reached about 5%, the tensile strength and elongation at break were over two times that of the pure PU. The thermal stability and the glass transition of the PU/OMMT nanocomposites also increased with increasing OMMT content. Jin et al. [45] investigation results on the viscoelasticity of PU-organoclay nanocomposites suggested that the addition of organoclay resulted in the increase in the elasticity of PU and the decrease in damping property, and significant improvement of the thermal stability of PU. However, it was also found that the addition of organoclay did not enhance the modulus of PU significantly, and that at high contents of soft segments the modulus of PU decreased with the addition of
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organoclay. Hysteresis results indicated that energy dissipation increased with increasing organoclay; at small deformation of 50 or 100%, the difference of dissipated energy between the two cycles was not affected by the organoclay. To the researchers, this meant that the part of the domain deformation energy was the same for the PU nanocomposites at different organoclay contents. It demonstrated that the hard domain deformation was not disturbed by the orientation of the organoclay platelets at small deformation. At the larger deformation of 200% a mutual influence between the domain deformation and the orientation of the organoclay platelets coexisted. Therefore, during the larger deformation the extra energy was required to overcome the effect. The more organoclay incorporated in the PU, the more the dissipated energy was. However, the researchers noted that the intercalation characteristic of organoclay by PU chains and the dispersion of organoclay in PU matrix could be key factors affecting the viscoelasticity of the composites. When inorganic fillers are dispersed in polymers, elastic modulus is typically increased with a concurrent sacrifice of the ultimate strength and elongation. Not surprisingly, the nanodispersed silicates resulted in a significant increase in modulus by a factor of two to three at 20 wt% organic layered silicate (OLS) content [28]. However, the ultimate strength remained comparable to that of the neat PU. These changes were also accompanied by retention of the native PU ductility and even an increase in strain to failure at relatively low OLS levels. Although in the bulk state silicates behaved as relatively rigid materials, TEM micrographs of nanocomposite morphology showed that individual silicate layers, because of their high aspect ratio and nanometer thickness, exhibited a measurable flexibility. Studies by Zhou et al. [46] found the Young’s modulus of PU coatings with nanosilica and with fumed silica was about the same, while the Young’s modulus of coatings with micron size silica was distinguishably smaller than both. The researchers attributed the findings to the reaction of isocyanate with hydroxyl groups on the surfaces of nanosilica resulting in a higher crosslinking degree for nanosilica embedded film than for micro-structured silica contained film. Another reason given was that there ought to have been greater interaction strength between nanosilica and organic matrix than micron size silica and organic matrix, since the former had greater specific surface area than the latter; fumed silica has to some extent similar particle characteristic to nanosilica. As concentration of microsilica increased to 10 wt%, the Young’s modulus of the coatings showed an apparent increase, especially under the high normal loads. It was concluded that the tensile strength and Young’s modulus were enhanced with the increasing content of nanosilica as displayed on Fig. 7; however, the elongation at break decreased as nanosilica content was increasing. Another report on PU/nanosilica composites showed a modulus of elasticity of 4.8 MPa while the elongation at break was measured to be 580% in comparison to that of pure PU; however, the elongation at break decreased as nanosilica content increased [46]. The UV absorbance in the wavelength of 290–400 nm increased as nano-SiO2 content increased. In contrast, for the PU coatings with fumed silica or micro-sized silica embedded, only hardness and abrasion resistance showed some
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Fig. 7 The tensile properties of PU films with different silica content (a) and Young’s modulus (b) of PU films containing different types of silica (reproduced from Ref. [56])
increase. In still developed experiment, nanotube sensors have been implemented in PU matrix for quantitative stress mapping in an elastic deformation situation [47]. This has been demonstrated by measuring the stress field in the vicinity of holes in polymer films; the experimental data were found to be in good agreement with the classical theory of Inglis as noted elsewhere [48]. The random dispersed nanotubes are more useful in practice because different stress components can be measured in the same sample. Further, Zhou et al. [38] prepared polyester-based PUs with embedded nanosilica particles. It was reported that addition of a small amount of nanosilica increased the hardness, abrasion resistance and tensile properties of the polymer films although these mechanical properties worsened at higher nanosilica contents. Also, according to the UV–Vis analysis it was noted that the absorbance and reflection of ultraviolet–visible light by the PU films increased as the nano-SiO2 content increased, especially at wavelengths of 290– 400 nm. The work further noted that the viscosity of the PU resin increased as the nanosilica content increased. Chiang et al. [49] report showed a significant advantages in the utilization of fullerenols as hyper-cross-linkers for the synthesis of PU nanoelastomers resulting in a substantial increase of polymer tensile strength for an order of magnitude from that of its linear polymer analog to a
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Fig. 8 Nominal stress-nominal strain curves for exfoliated second pseudogeneration HBP/Na+-MMTbased PU nanocomposites and intercalated second pseudo-generation HBP/ organically modified MMT composites for the overall compositions indicated (27°C and 4 mm/min for a 50-mm gauge length) (reproduced from Ref. [50])
maximum of 7.2 MPa. In Plummer et al. [50] report, the strength and stiffness in the exfoliated composites increased with respect to the neat matrix, however, there was a decrease in strain at break, as seen in on Fig. 8. The PU nanocomposites containing exfoliated OMS showed an increase in both strength and strain at break in comparison to nanosystems containing intercalated OMS. These results are in agreement with those of Tien and Wei [51], who reported a 34% increase in Young’s modulus. In mechanical properties of a novel segmented PU/clay nanocomposite based on PCL, MDI, butanediol, and PCL/clay prepolymer, about 1.4% PCL/clay in PU/clay resulted in a large increase in the elongation of PU/clay. However, when the amount of PCL/clay was 4.2%, the elongation of PU/clay was reduced drastically. This behaviour indicated that PU/clay could be transformed from an elastomer to a thermoplastic material as the amount of PCL/clay in PU/clay increased. Additionally, the lap shear stress of PU/clay was at least three times that of neat PU as a result of the influence of PCL/clay component [30]. A two fold increase in the tensile strength and a three fold increase in the elongation were found for PU composite containing 1% of benzidine-MMT (BZD-MMT) as compared to that of pure PU. Furthermore, both PU-based nanocomposites containing 1% of MMT modified with 12-aminolauric acid (12COOH-MMT/PU) and 1% BZD-MMT/PU exhibited lower water absorption properties than that of pure PU. In addition, both 1% 12COOH-MMT/PU and 1% BZD-MMT/PU exhibited lower water absorption properties than that of pure PU [31]. Scheme 1 shows the schematic drawing of proposed intermolecular interaction between PU molecules and (a) 12 aminolauric acid-MMT, (b) benzidine-MMT. Polymer sandwich composites, including PU-based composites, are widely used in aircraft engine nacelle, on wings for fuel tanks protection, tailplane panels for protection of stones and pebbles on take-off and landing, naval ships, human
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Scheme 1 The schematic drawing of possible intermolecular interaction between PU molecules and a 12 aminolauric acid-montmorillonite, b benzidine-montmorillonite (reproduced from Ref. [31])
vests and helmets for ballistic protection, automotive for collision and heat protection. In all these applications, the intrinsic properties of light weight and rigidity are used. These sandwiches, composed of a core of cellular material and two composite skins, are light (since their constituents are of low density), rigid in traction and compression (the composite materials have good mechanical properties) but also in bending since the foam core thickens the structure (and thus increases its quadratic moment while limiting its weight) and supports high bending moments. These properties are particularly interesting for producing functional structures that must sustain high stresses under normal conditions. During severe or extreme loadings (crashes or accidents), these structures must deform plastically and absorb the impact energy to protect either the rest of the structure or the vehicle occupants. A possible way of improving the properties of foam materials is through the inclusion of small amounts of nanoparticles (carbon nanotubes and nanofibres, TiO2, nanoclay, etc.) to improve the foam density and modulus properties. Up to now, MMT nanoclays have been the best candidate for foam reinforcement due to ease of processing, enhanced thermal–mechanical properties, wide availability and cost [52, 53]. Likewise, polyurethanes (PU) are core materials of choice due to their tailorable and versatile physical properties, ease of manufacture and their low costs. The use of epoxy resins filled with nanoparticles to construct either laminates or foams is relatively new. Moreover,
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the use of nanoparticles in such laminates [54], or foams in sandwich composite construction, is in its infancy but has been found to be both realistic and beneficial [4]. For instance, by using less than 5% by weight of nanoclay loadings, significant improvement in foam failure strength and energy absorption has been realised with over a 50% increase in the impact load carrying capacity when compared to a neat foam sandwich [55, 56]. Njuguna et al. [57] have recently conducted low velocity impact studies on nanophased polyurethane cores in sandwich structures. The investigation observed that nanophased sandwich structures are capable of taking higher peak loads than those made of neat polyurethane cores when subject to lowvelocity impact. It was found that the incorporation of MMT resulted in higher number of PU cells with smaller dimensions and higher anisotropy index (crosssections RI and RII). The obtained materials exhibited improved parameters in terms of thermal insulation properties. The results also show that nanophased sandwich structures are capable of withstanding higher peak loads than those made of neat polyurethane foam cores when subject to low-velocity impact despite lower density than that of neat PU foams. This is especially significant for multi-impact recurrences within the threshold loads and energies studied as shown on Fig. 9. Cao et al. [58, 59] claimed that nanoparticle loading caused significant changes in the cellular structures of the foam with cell dimension almost doubling with the inclusion of 3 wt% of TiO2 nanoparticles. On an average, the increment in strength and stiffness was 30 and 62%, respectively, over the neat system. Gain in strength
Fig. 9 Load versus time graph obtained from second impact test on sandwich structures using aluminium faceplates and PU/MMT nanofoams
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was attributed to the delay in the formation of initial cracks during loading. It was believed that nanoparticles embedded in the cell walls and edges and the associated interfaces surrounding the nanoparticles resisted crack formation/coalescence at the earlier stage of the loading. Accordingly, this allowed higher sustenance of load, but with the infusion of higher percentages of nanoparticles such as SiC, both thermal and mechanical properties began to deteriorate. During the flexural tests it was observed that failure always initiated on the tension side of the specimen. Cracks first appeared along the width of the specimen and then began to propagate through the thickness towards the compression side. Once the crack completed traversing the entire width and proceeded towards the thickness, the failure was very quick. Modes of crack propagation were almost identical with both neat and nanophased foams. However, there was a significant delay in the first appearance of the cracks in case of nanophased foams since they began to appear at a much higher load. It was proposed that nanoparticles embedded in the cell structures resisted the initial crack growth at the earlier stage of the loading that eventually contributed to higher failure load. SWNTs are considered to be the ideal reinforcing agents due to their exceptional mechanical properties, low density and high aspect ratio. Theoretical and experimental studies have shown that SWNTs have extremely high Young modulus, similar to that of graphite in-plane (*1,000 GPa). Of special interest is the study presented by Sen et al. [60], who fabricated SWNT-reinforced composite nanofibres and membranes using the electrospinning process. The membranes showed a nonlinear elastic behavior in the low stress region (0–2 MPa) and plastic deformation at higher stress. Additionally, the maximum stress at break was the tensile strength that increased for PU membranes on SWNT incorporation (Fig. 10). Further, compared to pure PU membranes, the tensile strength of as prepared PU-SWNT nanocomposites increased by 46% from 7.02 to 10.26 MPa—this enhancement in the mechanical properties was related to efficient load transfer to Fig. 10 Tensile stress–strain curves for PU membranes containing as-prepared (AP) and ester (EST) functionalized SWNTs. The SWNT-toPU weight ratio is 1:100 in the composite membranes (b, c) (reproduced from Ref. [53])
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the SWNTs in the composite material. On the other hand, nanocomposites fabricated with ester functionalized SWNT showed an increase of 104% in tensile strength when compared to electrospun pure PU. Better mechanical properties for the ester functionalized SWNT composites was attributed to improved dispersion of the SWNTs, but it was suggested that this could also have been a response of the polar functionalities in the modified SWNT to the opportunities offered by hydrogen bonding sites in the polymer matrix, or to reactions between free amine in the PU and the ester groups in the functionalized SWNTs. Generally, nanocomposites are materials whose components interact at a nanoscale level, a characteristic which confers great mechanical strength to these materials due to the quantum-scale interaction between these molecules which also gives it anti-fracture properties. The reason for this according to Griffin is that when a composite has fillers below a critical length, their strength even if cracked, is virtually equivalent to that of a solid crystal. This holds true for nanocomposites and explains the relative immunity of nanocomposites to fracture. Along this line of interest, one work [61] reported tests that showed significant improvements of impact strength of 27% for the phenolic resin/linear-PU system and 54% for the phenolic resin/C60-PU system, respectively, both with 3 wt% of linear-PU and C60-PU content. These results may be related to the intensive intermolecular hydrogen bonding that exists between phenolic resin and C60-PU as evident on FTIR data. Significant improvement in the toughness of the phenolic resin/C60-PU nanocomposite was also observed. However, Zilg et al. [29] reported both Shore A hardness and Young’s modulus, which reflected polymer stiffness, decreased slightly in modified PU/mica nanocomposites, whereas non-modified mica gave increased stiffness and only small increases of tensile strength with increasing filler content. In addition, the follow-up work [62] reported that interfacial coupling can be controlled in the nanostructure PU formation as a function of organophilic modification thus improving selected properties, such as toughness/ stiffness balance, heat distortion temperature and flame retardancy. Finnigan et al. [63] observed large improvements in stiffness on silicate addition, particularly for the soft PU elastomer of Shore Hardness 80A, which contained a higher fraction of soft segments. At a 7 wt% loading of organically modified layered silicate, a 3.2fold increase in Young’s modulus was noted. The improvement in stiffness was attributed to the good dispersion and delamination achieved, in addition to the strong interaction between PU matrix and OLS. Surprisingly, the addition of OLS resulted in a decrease in tensile strength and elongation; a fact worth further investigations. Nevertheless, the researchers attributed these effects to altered microphase morphology, excessive polymer-filler interaction, and an inhibition of the usual morphological changes that accompanied deformation in segmented PU, i.e. hard domain rotation, interchain slippage, fibrillation and soft segment crystallisation. The addition of layered silicates with high aspect ratios was observed to increase the hysteresis and permanent set of these PU elastomers. The enhancements of hardness properties in PU nanocomposites have been witnessed elsewhere too [46, 64]. In one report [41], this reinforcing effect was more evident when the elasticity of the PU matrix was increased at the temperature
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range above the Tg of hard segments and also in the situation where the deformation was large. Kuan et al. [65] work results on waterborne PU and polysilicic acid nanoparticles (PU/PSA nanocomposites) indicated that the wear index declined as the PSA particles content increased (below 15 wt%) over 500 and 1,000 wear cycles; the wear index was proportional to the reciprocal of the wearresistance. SEM revealed that the particles were well dispersed in the polymer matrix, suggesting that the nanocomposites exhibit good miscibility between organic and inorganic phases, so the wear resistance was enhanced by wearendurable PSA nanoparticles. The wear-resistance of the nanocomposites was greatest at 15 wt% PSA nanoparticle content. Adding polysilicic acid nanoparticles to waterborne PU increased the initial tensile modulus of the composites. The tensile modulus of the polysilicic acid nanoparticles/waterborne PU nanocomposite with 20 wt% silica increased from 20 to 41 MPa. Thus, the polysilicic acid nanoparticles play an important role in strengthening the composites by effectively transferring the stress between the silicon-containing nanoparticles and the PU matrix.
2.3 Polyaniline Nanocomposites Polyaniline (PANI) is a conducting polymer and its properties are strongly dependent on synthetic procedures, type of dopant, morphology, and other variables. PANI being electrically conducting in nature can be used for conductive adhesive, conductive ink, conductive paint, antistatic textile, and electrostatic discharge (ESD) materials. Applications for such composites are wide spread, these are used for interconnections, printed circuit boards, encapsulations, die attach, heat sinks, conducting adhesives, electromagnetic interference (EMI) shielding, ESD, and aerospace engineering. PANI-inorganic nanocomposites have also been proven to possess a wide range of properties such as electrical, mechanical, and structural properties because of synergistic effect owing to the intimate mixing between organic components in molecular level. The degradation behaviour of polyaniline nanocomposite is yet to be clearly understood and only a handful published work is available [66]. Liu et al. [67] synthesized Fe2O3 magnetic nanoparticles with size range of 50– 100 nm. The initial decomposition temperature of the composite coating was 250°C. The solar absorptivity (as) of the composite coatings was as high as 90%, and its emission rate (en) was reduced to 56%, while the relative efficiency of light–heat transition (as/en) was about 1.6. The nanocomposite coatings exhibited excellent abrasion resistance, weatherability and water resistance due to the formation of a three-dimensional network structure during the thermal curing process. The results indicate that the nanocomposite material could be used as a solar light– heat transition coating that could be employed in solar hot-water collection. For photodegradable packaging materials, nano-TiO2 can be used. The PANI–TiO2 nanocomposite powders showed highly enhanced photodegradation and the
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photodegradation increased with decreasing ratios of PANI:TiO2. A weight loss of about 6.8% was found for the PANI–TiO2 (1:3) nanocomposite; however, the weight loss of the PANI–HCl powder was only 0.3% after being irradiated for 60 h under air. The photocatalytic degradation of the nanocomposite powders accompanied the peak intensity decrease in the FTIR spectra at 1,235 cm-1, attributed to C–N stretching mode for benzenoid unit, and the depigmentation of the powders due to the visible light scattering from growing cavities. The elemental analysis and X-ray photon spectroscopy (XPS) analysis of the composite showed that the bulk and surface concentration of N decreased with irradiation [68]. Lanthanum (La)doped Fe3O4 magnetic nanoparticles were prepared in aqueous solution at room temperature, then La-doped Fe3O4–polyaniline (PANI) nanocomposites containing a dispersion of La-doped Fe3O4 nanoparticles were synthesized via in situ polymerization of aniline monomer. The La-doped Fe3O4–PANI composite presented core–shell structures; polyaniline covered the La-doped Fe3O4 completely. The specific saturated magnetization of La-doped Fe3O4–PANI depended on the starting material of La-doped Fe3O4. Lee and Char [69] have found that the PANI/Na-MMT nanocomposites were more thermally stable than the physical mixture of PANI and Na-MMT. Polyaniline undergoes a three-step thermal decomposition. The weight loss in the third step, which is attributed to polyaniline backbone decomposition, was found to be maximum at 530°C for pure PANI and this was shifted 25°C more for PANI/NaMMT nanocomposite. From the XRD investigation after TG analysis, it can be concluded that the PANI chains residing outside the silicate layers decomposed mostly, whereby chains residing inside the layers to a small extent. So the shielding effect of intercalation into the layers imparts the thermal stability to polymeric materials.
2.4 Poly(Ethylene Terephthalate) Nanocomposites Polyethylene terephthalate (PET) exists both as an amorphous (transparent) and a semi-crystalline (opaque and white) thermoplastic, and can be made into either as resin, fibre or film. The semi-crystalline PET has good strength, ductility, stiffness and hardness while the amorphous PET has better ductility. PETs are well suited for flexible laminates, customized for specific performance parameters, are of interest to aerospace programs such as lighter-than-air vehicles, balloon systems, decelerator systems, flexible inflatable structures and pressure vessels, inflatable structures and pressure vessels. The material requirements for these applications include high strength-to-weight ratio and modulus, low gas permeability, pressure retention and the capability to survive in harsh atmospheric, marine and/or stratospheric environments for extended periods of time. These flexible laminates achieve a significant weight savings over woven fabrics of similar strengths by eliminating strength and modulus loss and other structural deficiencies caused by
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crimping of yarns during the weaving process. The absence of crimp in non-woven fabrics results in a linear elastic response that allows for ease in predicting material properties and simplification of structural models. These flexible composites afford the ability to specify structural properties, oriented to meet any design requirement. Parts can be manufactured with complex. Zeng et al. [70] prepared poly(ethylene terephthalate) (PET)-based nanocomposite fibres by melt spinning three types of PET/polyhedral oligomeric silsesquioxane (POSS) composites. These composites were made by either melt blending POSS with PET at 5 wt% loading level (non-reactive POSS and silanol POSS) or by in situ polymerization with 2.5 wt% reactive POSS. Significant increases in tensile modulus and tensile strengths were achieved in PET fibres with non-reactive POSS at room temperature. The high temperature modulus retention was found to be much better for PET/silanol POSS fibre when compared to that of control PET. Although other PET/POSS nanocomposite fibres tested did not show this high retention of modulus at elevated temperatures, PET/isooctylPOSS nanocomposite fibres did show increased modulus at elevated temperature compared to that of PET. Higher compressive strengths, compared to PET fibres, were observed for all three nanocomposite fibres. Gel permeation chromatography measurement suggested that there was no significant change in molecular weight during preparation of PET/POSS nanocomposites. SEM observations suggested that there was no obvious phase separation in any of the three PET/POSS systems. The fibre spinning and mechanical performance with 10 and 20 wt% of trisilanolisooctyl POSS were also investigated. It was noted that the nanocomposites with higher concentrations of this nanofiller can be spun without any difficulty. At room temperature, the fibre tensile modulus increased steadily with the POSS concentration while fibre tensile strength showed no significant change. The elongation at break decreased significantly with increasing of POSS concentration. The high-temperature moduli of PET/POSS nanocomposite fibres were found to be rather variable, likely due to the modest compatibility between filler and polymers, which lead to structural anisotropy within the composite.
2.5 Polyimide Nanocomposites Fiber reinforced PMR polyimides are finding increased acceptance as engineering materials for high performance structural applications. Prepreg materials based on this novel class of highly processable, high temperature resistant polyimides [71]. Applications of polyimide foams in aerospace vehicles are include thermal-insulating systems for cryogenic tank, honeycomb sandwich structures, radomes and chairs. Ogasawara et al. [72] directed their investigations toward improvement of heat resistance of phenylethynyl terminated imide oligomer (Tri-A PI) by loading of MWNT. They fabricated the Tri-A PI/MWNT nanocomposites containing 0, 3.3, 7.7, and 14.3 wt% MWNT using a mechanical blender without any solution (dry condition) for several minutes. The volume fraction of MWNT were
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Fig. 11 Effect of the MWCNT concentration on Young’s modulus of the composites (reproduced from Ref. [65])
calculated to be 2.3, 5.4, 10.3 vol% from the density of the MWNT (1.9 g/cm3) and the cured polyimide (1.3 g/cm3). Scanning electron micrographs showed the particle size of the imide oligomers to be in the range of 0.1–10 lm, and MWNT were not dispersed uniformly in the mixture. The loss of aspect ratio during the mechanical blending was not significant, therefore the MWNT were flexible for mechanical blend process with the imide oligomers. The preparation of the nanocomposite involved the melt mixing of imide oligomer/MWNT at 320°C for 10 min on a steel plate in a hot press, and then curing at 370°C for 1 h under 0.2 MPa of pressure with PTFE spacer (thickness 1 mm). Tensile tests on the composites showed an increase in the elastic modulus and the yield strength, and decrease in the failure strain. Figure 11 shows the effect of the MWNT concentration on Young’s modulus of the composites. Dynamic mechanical analysis showed an increase in the glass transition temperature with incorporation of the carbon nanotubes. The experimental results suggested that the carbon nanotubes were acting as macroscopic crosslinks, and were further immobilizing the polyimide chains at elevated temperature. As to the reason why dispersed MWNT increased the heat distortion temperature, the researchers explained that the dispersed MWNT impedes the molecular motion in polyimide network at elevated temperature. The other property improvements in this material are that MWNT showed some potential for controlling electric conductivity and electro-magnetic wave absorbability. Although static properties were obtained, discussions were not given, and it is evident that more research work would be required to prove that the suggested phenomenon is a true cause of higher glass transition temperature.
2.6 Polyarylacetylene Nanocomposites Polyarylacetylene (PAA) is going through increasing development in the field of advanced heat resistant composites owing to its outstanding heat resistance and excellent ablative properties. PAA is developed increasingly as the matrix for high
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temperature composites of next generation in carbon–carbon composites, ablation materials and structural materials owing to its outstanding heat resistance and excellent process properties. Resulting mainly from the diethynylbenzene (DEB), PAA is a highly cross-linked aromatic polymer that contains only carbon and hydrogen when it is cured by means of addition polymerization, which makes it own excellent thermal stability and oxidation. When the PAA is heated to high temperatures in an inert environment, only about 10 wt% is volatilized, while the remaining 90 wt% is carbon char, which means far less volatile material is generated and minimal shrinkage is associated with pyrolysis. Furthermore, PAA resin is liquid or solvable and fusible solid at room temperature so that it provides better processing flexibility, being applicable for conventional curing processes like compression molding process, vacuum bagging method, resin transfer molding. However, the bad wettability between carbon fibre and PAA resin from the nonpolar structure of PAA and chemical inertness of carbon fiber causes the weak interfacial adhesion between fiber and non-polar PAA resin. The main potential applications of PAA resin are in conventional resin matrix composites with ultra-low moisture outgassing characteristics and improved dimensional stability suitable for spacecraft structures, as an ablative insulator for solid rocket motors, and as a precursor for carbon–carbon composites. Carbon fibre reinforced PAA composites (carbon fibre/PAA) undoubtedly play a very important role in all these fields. Unfortunately, the mechanical properties of the carbon fibre/PAA material are not yet sufficiently satisfactory to replace the widely used heat resistant composites such as carbon or graphite reinforced phenolic resins. The mechanical properties of carbon fibre reinforced resin matrix composites depend on the properties of carbon fibre and matrix, especially on the effectiveness of the interfacial adhesion between carbon fibre and matrix. Polyarylacetylene has high content of benzene rings and hence a highly crosslinked network structure, which render the material brittle. Moreover, the chemical inert characteristics of the carbon fibre surface lead to weak interfacial adhesion between fibres and non-polar PAA resin. To ensure that the material could be used safely in complicated environmental conditions and to exploit the excellent heat resistant and ablative properties more effectively, it is necessary to improve the mechanical properties of the carbon fibre/PAA composites. To achieve this purpose, two kinds of methods can be used. One method is to improve the properties of PAA resin by structural modification or by co-mixing other resins, such as phenolic resin. The other is treatment of carbon fibre surface. The latter method has been studied for a long time and several methods, such as heat treatment, wet chemical or electrochemical oxidation, plasma treatment, gas-phase oxidation, and high-energy radiation techniques have been demonstrated to be effective in the modification of the mechanical interfacial properties of advanced composites for engineering applications. In Zhang et al. [73] investigations, for instance, carbon fibres were subjected to oxidation–reduction processes followed by vinyltrimethoxysilanes–silsesquioxane (VMS–SSO) preparation method to improve the interfacial mechanical properties of the carbon fibre/PAA composites. The carbon fibre surface treatment process is shown in Fig. 12.
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Fig. 12 Schematic diagrams of carbon fibre treatment process (sample 1 is oxygen plasma oxidation, sample 2 is LiAlH4 reduction, sample 3 is VMS–SSO coating) (reproduced from Ref. [66])
Polar functional groups, including carboxyl and hydroxyl, were implemented on carbon fibre surface after the oxygen plasma oxidation treatment. The quantity of carboxyl groups on carbon fibre surface was decreased and that of hydroxyl moieties on carbon fibre surface was increased after the LiAlH4 reduction treatment [74]. The VMS–SSO coating was grafted onto the carbon fibre surface by the reaction of the hydroxyl groups in VMS–SSO and that on carbon fibre surface. The VMS–SSO coating concentrations and treatment time were decided according to Zhang et al. [75] who had optimized VMS–SSO coating treatment parameters. The investigation found out that interlaminar shear strength of the PAA/carbon fibre composites was increased by 59.3% at the end of treatment [73]. This kind of modification method could be widely used in different resin matrix composites by changing the functional groups on silsesquioxanes according to that on the resin.
2.7 Poly(ether ether ketone) PEEK Nanocomposites Poly(aryl–ether–ether–ketone) (PEEK) is a high performance semicrystalline thermoplastic with wide applications in aerospace, automotive, coating, electrical insulating material and oil fields. Highly crystalline PEEK materials display excellent mechanical, thermal and chemical resistance properties which allow filled or unfilled PEEK to be used in various hostile environments. Even though PEEK materials have been shown to perform satisfactorily up to temperatures of 200°C, there are a few current applications that require materials with even higher upper temperature limits which are targets for improvement using nanofillers. Jen et al. [76] manufactured PEEK/silica nano-composite laminates and also studied their mechanical responses. The experimental procedure were as follows: firstly, the nanoparticles were diluted in alcohol (50 ml alcohol:2 g SiO2) and stirred uniformly, then 16 plies of [0/90]4s cross-ply and [0/±45/90]2s quasi-isotropic prepregs were cut, SiO2 solution was then spread on the prepreg in a temperature-controlled box, and later weighed the nanoparticles after evaporation
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Fig. 13 Pressure–temperature profile of the curing process of AS-4/PEEK APC-2 nanocomposites (reproduced from Ref. [76])
of alcohol in the range of 111–148 mg/ply. A repeat on spreading for 5, 8, 10, 15 plies was the next step followed by curing (the curing process is shown in Fig. 13) the stacked plies in a hot press to form a laminate of 2 mm thick. Next, the laminates were cut into specimens and tested according to ASTM D3039M. The tensile tests were repeated at 50, 75, 100, 125, 150°C to receive respective stress–strain curves, strength and stiffness, and the obtained data compared with the original PEEK laminate (no SiO2 nanoparticles) to find the optimal SiO2 content. From tensile tests it was found out that the optimal content of nanoparticles (SiO2) was 1% by total weight. The ultimate strength increased by about 12.48% and elastic modulus by 19.93% in quasi-isotropic nano-laminates, whilst, the improvement of cross-ply nano-composite laminates was less than that of quasi-isotropic laminates. At elevated temperatures the ultimate strength decreased slightly below 75°C and the elastic modulus reduced slightly below 125°C, however, both properties degraded highly at 150°C (&Tg) for the two layups. Finally, after the constant stress amplitude tension–tension (T–T) cyclic testing, it was found that both the stress-cycles (S–N) curves were very close below 104 cycles for cross-ply laminates with or without nanoparticles, and the S–N curve of nano-laminates slightly bent down after 105 cycles. Sandler et al. [77] produced poly(ether ether ketone) nanocomposites containing vapour-grown CNF using standard polymer processing techniques. Macroscopic PEEK nanocomposite master batches containing up to 15 wt% vapour grown CNF were prepared using a co-rotating twin-screw extruder with a length-to-diameter ratio of 33. The processing temperatures were set to about 380°C. The strand leaving the extruder was quenched in a water bath, air dried and then regranulated followed by drying at 150°C for 4 h. Tensile bars according to the ISO 179A standard were manufactured on an injection moulding machine at
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processing temperatures of 390°C, with the mould temperature set to 150°C. Prior to mechanical testing, all samples were heat treated at 200°C for 30 min followed by 4 h at 220°C in an attempt to ensure a similar degree of crystallinity of the polymer matrix. Macroscopic tensile tests were performed at room temperature— the cross-head speed was set to 0.5 mm/min in the 0–0.25% strain range and was then increased to 10 mm/min until specimen fracture occurred. Evaluation of the mechanical composite properties revealed a linear increase in tensile stiffness and strength with nanofibre loading fractions up to 15 wt% while matrix ductility was maintained up to 10 wt%. Electron microscopy confirmed the homogeneous dispersion and alignment of nanofibres. An interpretation of the composite performance by short-fibre theory resulted in rather low intrinsic stiffness properties of the vapour-grown carbon nanofibre. DSC showed that an interaction between matrix and the nanoscale filler could occur during processing. However, such changes in polymer morphology due to the presence of nanoscale filler need to be considered when evaluating the mechanical properties of such nanocomposites. Schmidt [78] investigation involved multifunctional inorganic–organic composite sol–gel coatings for glass surfaces. The sol–gel process allowed the fabrication of ceramic colloidal particles in the presence of organo alkoxy silanes carrying various functions and the synthesis of multi-functional transparent inorganic–organic composites. The report claimed that, in addition, these composites can be used as controlled release systems or designed as gradient systems. Using this approach, a coating with a very low surface free energy (antisoiling properties) and temperature stability up to 350°C, a controlled release system for permanent wettability (anti-fogging) and systems containing metal colloids for optical effects were developed. Lin [79] and Wang et al. [80] studied the effect on wear and friction by adding SiC nanoparticles in PEEK. The latter studied the effect of the synergism between nanometer SiC and PTFE on the wear of PEEK. The PEEK fine powders (ICI grade 450P, g = 0.62) in a diameter of approximately 100 lm, were prepared. The nanometer SiC used as filler had the size smaller than 80 nm. The PTFE powders (diameter 25 lm), nanometer SiC and PEEK were fully mixed ultrasonically, dispersed in alcohol for *15 min. Then the mixture was dried at 110°C for 6 h to remove the alcohol and moisture. Finally, the mixture was moulded into the block specimens by compression moulding, in which the mixture was heated at a rate of 10°C min-1 to 340°C, held there for 8 min, and then cooled in the mould to 100°C. After releasing from the mould, the resultant block specimen was prepared for friction and wear tests. A tribological study found that the incorporation of PTFE into 3.3 vol% nanometer SiC filled PEEK had a detrimental effect on the tribological properties of SiC–PTFE–PEEK composite. The morphologies of worn surfaces and the properties of transfer films deteriorated, while the load-carrying capacity of the SiC–PTFE–PEEK composite was also adversely affected. The researchers claimed the reason for this was due to SiFx, which was formed on the original surface and worn surface during the compression moulding process and sliding friction process as a result of the chemical reaction between nanometer SiC and PTFE. The chemical reaction and the formation of SiFx dominated the tribological behaviour of the SiC–PTFE–PEEK
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composites filled with various contents of PTFE and 3.3 vol% nanometeric SiC. When the PTFE volume percentage was low, then the SiFx caused the friction and wear of the SiC–PTFE–PEEK composite to rise. However, at high volume percents the low friction PTFE dominated the friction and wear behaviours and the friction decreased as the percentage of PTFE increased. The chemical reaction and the formation of SiFx led to changes in the worn surface morphologies and detrimental effect on the characteristics of the transfer films.
2.8 Poly(p-phenylene benzbisoxazole) (PBO) Nanocomposites Poly(p-phenylene benzbisoxazole), a rigid-rod polymer, is characterized by high tensile strength, high stiffness, and high thermal stability. Kumar et al. [81] found out that PBO/CNT reinforced fibres exhibited twice the energy absorbing capability than the plain PBO fibres. The nanocomposites were prepared as follows: into a 250 ml glass flask, equipped with a mechanical stirrer and a nitrogen inlet/ outlet, were placed *4.3 g (0.02 mol) of 1,4-diaminoresorcinol dihydrochloride, *4 g (0.02 mol) of terephthaloyl chloride, and *12 g of phosphoric acid (85%). The resulting mixture was dehydrochlorinated under a nitrogen atmosphere at 65°C for 16 h and subsequently at 80°C for 4 h. At this stage, 0.234 g of purified and vacuum-dried HiPco nanotubes was added to the reaction flask. The mixture was heated to 100°C for 16 h while stirring and then cooled to room temperature. P2O5 (8.04 g) was added to the mixture to generate poly(phosphoric acid) (77% P2O5). The mixture was stirred for 2 h at 80°C and then cooled to room temperature. Further P2O5 (7.15 g) was then added to the mixture to bring the P2O5 concentration to 83% and the polymer concentration to 14 wt%. The mixture was heated at 160°C for 16 h with constant stirring. Stir opalescence was observed during this step. The mixture was finally heated to 190°C for an additional 4 h while stirring. An aliquot of the polymer solution was precipitated, washed in water, and dried under vacuum at 100°C for 24 h. An intrinsic viscosity of 14 dl/g was determined in methanesulfonic acid at 30°C. A control polymerization of pure PBO was also carried out under the same conditions without adding SWNT. For PBO/SWNT (90/10) composition, 0.47 g of purified HiPco tubes (SWNT) was added to the mixture. The sequence of steps and polymerization conditions remained the same as those for PBO/SWNT (95/5) composition. Intrinsic viscosity values of PBO and PBO/SWNT (90/10) were 12 and 14 dl/g, respectively. Singlewalled nanotubes were well dispersed during PBO synthesis in PPA. PBO/SWNT composite fibres were successfully spun from the liquid crystalline solutions using dry-jet wet spinning. The addition of 10 wt% SWNT increased PBO fibre tensile strength by about 50% and reduced shrinkage and high-temperature creep. The existence of SWNT in the spun PBO/SWNT fibres was evidenced by the 1,590 cm-1 Raman absorption band.
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3 Concluding Remarks Advancements in the nanotechnology industry promise to offer improvements in capabilities across a spectrum of applications. This is of immense strategic importance to the aerospace sector which has historically leveraged technological advances. In particular, nanomaterials represent a class of exciting new materials with extraordinary electrical, thermal, and mechanical properties. Taking advantage of these intrinsic properties relies on the ability to systematically synthesize, characterize, and integrate standardized materials into actual commercial products. Appropriate characterization of the properties (diameter, length, purity, surface area, bundling, etc.) has been highlighted with microscopic, spectroscopic, thermal, and surface area analysis, showing the importance of these techniques in moving towards material utilization. A few points can be emphasized here: (a) nanocomposites differ from conventional composites in that the mixing of phases occurs over a much smaller length scale in comparison to the micrometer length scale of conventional composites, resulting into nanocomposite materials with properties substantially different from those of the parent end members, (b) proper engagement of the nanoparticles in composite systems depends strongly on the ability to homogeneously disperse them throughout the matrix without destroying their integrity, and (c) the properties of the matrix, the distribution and properties of the filler as well as the nature of their interface control the behaviour of a typical composite material. In these terms, the nanoparticles often strongly influence the properties of the composites at very low volume fractions. This is mainly due to their small interparticle distances and the conversion of a large fraction of the polymer matrix near their surfaces into an interphase of different properties in addition to the consequent change in morphology. As a result, the desired properties are usually reached at low filler volume fraction, which allows the nanocomposites to retain the macroscopic homogeneity and low density of the polymer. Besides, the geometrical shape of the particles plays an important role in determining the properties of the composites. Consequently, nanocomposites have attracted much scientific and industrial interest in recent times. It is emphasized that polymers are subjected to destructive factors such as mechanical stress, the presence of different chemicals, UV-light, ablation and high temperatures throughout shelf- and service-life. These factors cause degradation and ultimately affect performance and lifetime of the polymers that are sometimes stored for long periods of time. Therefore it is important to know how long and under which conditions the polymers may best be stored with minimum deterioration of the properties. According to the lifetime stages of polymers, the relevant processes are classified as melt degradation; long-term heat ageing and weathering based on the mechanisms involved, i.e. thermomechanical, thermal, catalytic and radiation-induced oxidations and environmental biodegradation. The commercial importance of polymers has been driving intense applications in the form of
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composites in aerospace needs. Performance during use is a key feature of any composite material, which decides the real fate of products during use in aerospace indoor and outdoor applications. Whatever the application, there is often a natural concern regarding the durability of polymeric materials partly because of their useful lifetime, maintenance and replacement. The deterioration of these materials depends on the duration and the extent of interaction with the environment. Intense research is required in this area to allow real nano enhanced products applications.
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Friction and Wear of Rubber Nanocomposites Containing Layered Silicates and Carbon Nanotubes D. Felhös and J. Karger-Kocsis
Abstract This chapter gives a survey on the tribological performance of organoclay and carbon nanotube reinforced rubbers of both conventional (thermoset) and thermoplastic versions. The unlubricated friction and wear of rubbers were grouped in abrasion-, sliding- and rolling-types in order to support the overview. It was highlighted that the coefficient of friction and specific wear rate strongly depend on the configuration and testing parameters of the tribotests used. It was demonstrated that the incorporation of the above nanofillers is not always associated with improved resistance to wear and reduced coefficient of friction. Further experimental studies, data mining through proper statistical techniques, and extensive modeling works are needed to realize the potential of the above nanofillers in tribological applications, and to deduce relationships between wear and other characteristics (e.g. network-, mechanical response-related) of rubbers.
D. Felhös (&) Department of Polymers Engineering, Faculty of Materials Science and Engineering, University of Miskolc, 3515 Miskolc, Hungary e-mail:
[email protected] J. Karger-Kocsis Department of Polymer Technology, Faculty of Engineering and Built Environment, Tshwane University of Technology, Pretoria 0001, Republic of South Africa J. Karger-Kocsis Department of Polymer Engineering, Faculty of Mechanical Engineering, Budapest University of Technology and Economics, 1111 Budapest, Hungary
V. Mittal et al. (eds.), Recent Advances in Elastomeric Nanocomposites, Advanced Structured Materials, 9, DOI: 10.1007/978-3-642-15787-5_13, Springer-Verlag Berlin Heidelberg 2011
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1 Introduction 1.1 Nanotubes Nowadays considerable research interest is focused on carbon nanotubes (CNT) and their use in structural and functional polymer composites. CNT-reinforced elastomers are no exceptions in this respect. Already a large body of work is devoted to investigate the applicability of this novel filler in various rubber systems. The most exciting properties of CNTs are their high strength and extremely high modulus (*1 TPa), excellent electrical and thermal conductivities, high specific surface area and very high aspect ratio (ratio of length to diameter) ([1] and references therein). These properties of CNTs are very promising also for rubber products. The ongoing research and development works are dealing with property improvements and generation of novel functional properties. The commonly used CNT types are single-walled (SWCNT) and multi-walled (MWCNT) versions. However, one can find more special types like radial singlewalled and nanohorn-like CNTs. The dispersion quality of the nano-sized fillers in the matrix material controls the properties of the corresponding nanocomposites. The larger is the active filler surface in contact with the matrix, the larger the reinforcing effect is. It is obvious that avoiding the agglomeration of nanofillers with proper dispersion techniques has paramount importance. Inside the agglomerates there are not enough polymer matrix molecules to guarantee the stress transfer towards the reinforcing nanofillers. As a consequence the filler cannot bear the mechanical load. In addition, the non-wetted inner part of the agglomerate has stress concentration effect and thus it causes premature failure during loading. Moreover, good nanofiller dispersion is more economical as the same filler-matrix contact surface is reached with less filler. Good dispersion of the nanotubes can be reached by applying dispersion-, solution- or melt-blending methods. The greatest challenge is to overcome the van der Waals forces between the CNTs in order to disintegrate their agglomerates. To decrease the van der Waals interactions between the nanotubes, the CNTs are usually surface treated for example in acid bath, adding coupling agents, or applying plasma-, thermal- and laser-treatments [2–6]. The agglomerates can be broken in suitable dispersion media (aqueous solutions, solvents) mechanically. This occurs by ball milling [7, 8], sonication [9–12], or applying high rotation speeds in a dissolver [13, 14]. The surface treated and deagglomerated nanotubes can be mixed with the rubber in a kneader or on an open mill. It was observed, that the common use carbon black (CB) or silica with CNT helps to disperse the nanotubes homogeneously [15–17]. Other possibility is the in situ polymerization of elastomer in presence of CNT [18–20] which results in a homogeneous dispersion of the latter. Similar to CB, CNTs are also active fillers. This means that their incorporation in elastomers leads to an increase in the ‘‘apparent’’ crosslink density. As a consequence, the hardness, stiffness and strength of the rubbers increase with
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increasing filler content [21–28]. On the other hand, when CNTs are incorporated in large amount in the rubber matrix, they easily reagglomerate. Though the stiffness and the strength increase with filler loading, even in case of reagglomeration, the ultimate strain to failure of the related rubber nanocomposite is strongly reduced [11, 24, 25, 33]. Similar results are obtained when CNTs are added in thermoplastic elastomers [21, 33]. This scenario may change, however, when semicrystalline polymer phase is present in the thermoplastic elastomer, as in PP/EPDM blends [29]. In this case the CNTs are incorporated in the thermoplastic phase, influencing heavily the crystallization process of the thermoplast. Because of this feature, only a small amount of CNT filler should be introduced (\0.5%) to improve the mechanical properties of the corresponding thermoplastic elastomer. Note that similar conclusion was deduced in a study in which carbon nanofiber was mixed with an olefinic thermoplastic dynamic vulcanizate [30]. Orientation of the nanotubes in tension direction was observed during uniaxial stretching of CNT reinforced elastomers [24]. Similar to traditional composites, the orientation of the nanotubes before curing influences largely the mechanical properties of the of the CNT-elastomer nanocomposites [1]. Swelling tests showed similar results for CNT-filled rubbers, namely the equilibrium swollen volume of the nanocomposites decreased with increasing filler loading, confirming a strong interaction between the filler and the matrix [13, 31]. Dynamic-mechanical thermal analysis (DMTA) indicated that the maximum of the mechanical loss (tan d) peak decreases with increasing CNT filler content. Interestingly it is not clear yet, how the incorporated CNT affect the glass transition temperature (Tg). In the literature contradictory results are reported related this question. Slight increments or unchanged Tg values were reported in Refs. [7, 8, 26, 28, 31]. On the other hand, a significant Tg decrease was observed in Ref. [25]. Further open questions are linked with the explanation of the widely investigated Mullins and Payne effects [24, 25]. The former may be caused by the relaxation of the molecule chains, or through molecule-detachments from the filler surface. The Payne effect is commonly explained through the breakdown of the filler particle agglomerates due to large strains imposed on the material [16]. For the producers of rubber goods it is of great interest how the curing parameters are affected by the CNT filler. It turned out that the necessary vulcanization temperature increases with the incorporation of CNTs [7, 26, 28], and parallel to that the curing time decreases [8]. The electrical conductivity of rubber compounds is an important feature for tire compounds. The CB containing rubber compounds have relatively small electric resistance, but applying novel fillers such as silica changes this situation [32]. Discharge phenomena after static charging are inconvenient for the passengers and may damage the sensitive electronic equipments of modern cars, as well. To avoid static charging the electric conductivity of polymeric materials should be enhanced. General observation is that incorporation of CNTs increases the conductivity both of elastomers [16, 17, 33] and thermoplastics [34, 35]. The rapid increase in the electric conductivity due to increasing CNT concentration is
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explained by the aspect ratio of the nanotubes ([100), which is higher than that of short fibers (\30) or of CB (*1) [36]. Due to the high aspect ratio of CNTs more electrical contacts are created in the nanocomposites and the electric percolation threshold reached at very low filler loading. Similar to the mechanical properties of fiber reinforced composites, Kim et al. [1] reported, that the orientation of the CNTs enhances also the electrical properties of the nanotube-filled elastomer composite.
1.2 Organoclays Layered silicates in their pristine (crude) form are no active fillers for polymeric materials. Their special layered structure can be exploited for reinforcement only after chemical treatment. During this treatment the hydrophilic alkali metal cations at the surface of the silicate layers are replaced by bulky organophilic cations bearing usually long apolar chains. Through this cation exchange the sheet-to-sheet (interlayer) distance increases from the original ca. 1.5 nm to ca. 2.5 nm in the corresponding organoclays (OCs) [37]. This organophilic treatment renders the silicate layers hydrophobic and the interlayer van der Walls forces are dramatically decreased. This supports the intrusion of the rubber molecules between the layers during in situ polymerization [38–40], or melt mixing [41–43]. Depending on the compatibility of the organoclay with the polymer molecules, an intercalated or fully exfoliated clay structure is the final outcome. The modulus and the strength of the nanocomposites can highly be increased when the clay layers are fully exfoliated (delaminated) in the elastomer matrix [39, 41]. The larger is the distance between the modified clay layers initially the easier the generation of a well dispersed nanocomposite structure is [37]. Due to this, the physical properties of the nanocomposites are largely influenced trough the treatment of the nanoparticles [41, 44, 45]. In contrast to intercalated structures, the stiffness and elongation to break values increase monotonously with filler content for fully exfoliated nanocomposites [41]. Similar observations were reported by Yang [46], who pinpointed also the change in the viscoelastic properties between exfoliated and intercalated nanocomposites. The excellent properties of nanocomposites with exfoliated clay are underlined by Yousoh [39]. On the other hand, several studies suggested that the OC layers are energetically more stable when they are ranged close to each other. This can lead to de-exfoliation or de-intercalation processes of the exfoliated or intercalated structures during the relaxation of the polymer molecules [41, 47]. Contrary to increasing tensile strength of elastomeric organoclay (OC) nanocomposites, decreasing tear strength [43] and decreasing fracture toughness [48] were observed for epoxidized natural rubber (ENR)/OC and for polyamide (PA-6)/ maleinated polyolefin elastomer/OC composites, respectively, with increasing OC content.
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Special properties of the elastomer/OC nanocomposites are the following: increased thermal stability [38, 40, 43, 49–53], increased flame retardancy [53, 54], improved barrier properties to gases and liquids [40, 51, 53, 55]. The increased thermal stability due to organoclay incorporation, however, did not affect the recyclability via remelting as shown on example of thermoplastic elastomers. This was traced to the fact that the nanofiller was located in the thermoplastic phase [56]. In many applications the optical properties of the elastomers are important. Wang [42] reported that the OC particles tend to agglomerate with their increasing content. This decreases the transparency of the initially transparent elastomeric matrix. Other studies suggested that the critical percolation filler content can be detected by considering changes in the linear viscoelastic material properties [57] or in the initial modulus of composites [58].
2 Friction and Wear of Elastomers 2.1 Abrasion-Type Dominant part of the tribological investigations on elastomers was done for the car tire industry. Cardinal questions in this respect are the wear and frictional properties of the tires. Note that these properties are highly valued as they affect the safety of the car passengers, as well. Because of abrasive wear controls the performance of car tires, this is the most widely studied wear phenomenon. To compare the abrasion resistance of different rubber mixtures, without field trials, several standard laboratory abrasive test methods were developed, such as the Taber Abrader method [59]. Simplified laboratory test methods provide only comparative test results on the investigated rubber goods. To observe and analyze the wear mechanisms under abrasive conditions specific methods were developed. In the corresponding tests a razor blade or a needle abrade the rubber surface, simulating a single roughness asperity of the abrasive surface. Detailed description of the latter test methods is disclosed in Refs. [74–77; blade on rubber plate test or blade abrader], and in Refs. [69, 78, 79; scratch or needle on rubber plate test]. Applying these tests one can mimic and investigate the effect of one single roughness peak of an abrader. The dimensions of the rubber tongues, angles of the crack propagation, crack density, etc… were determined and different mechanical models proposed to explain the abrasion properties of the investigated rubbers [67, 74, 82–84]. Generally the abrasion is a very harsh action which produces a rough surface. For the heavy surface destruction hardly any schematic interpretation can be given. A typical unidirectional abraded surface structure is depicted in Fig. 1 showing abrasive grooves along with wear debris. This kind of wear can be ‘‘produced’’ also by using the laboratory abrasion test rig shown in Fig. 2.
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Fig. 1 Unidirectional abraded surface of conventional, thermoset elastomers
Fig. 2 Abrasive Pin on Plate (POP) tests for TPO based thermoplastic elastomer nanocomposites. Note: the wear of the specimens was collated based on the diameter of the wear tracks
2.1.1 Conventional Rubbers Extended investigations were performed by Faulkner et al. [15] who studied the mechanical and tribological properties of CNT reinforced fluororubber (FKM) and hydrogenated acrylonitrile–butadiene (HNBR). The deduced abrasive wear mechanisms suggested that the abrasion resistance of the investigated rubbers depends on their tear strength directly. However, the increase of the tear strength was not accompanied by a monotonous increase in the abrasion resistance. This means that also other mechanical parameters play an important role in the wear resistance of the material. For bisphenol cured FKM higher abrasion resistance was noticed at 10 phr CNT than at 30 phr N-990 CB content (DIN abrasion test (ASTM D5963-04) results are 11.2 and 18.1 mg, respectively). In this case the change in the tear strengths was also remarkable: 55.3 and 18.4 kN/m for the 10 phr CNT and 30 phr N-990 CB fillers, respectively. Slightly different abrasive wear resistance was found for the FKM when instead of bisphenol, peroxide curing was used. In this case the DIN abrasion tests resulted in 7.5 and 7.7 mg material losses for 10 phr CNT and 35 phr N-990 CB contents, respectively. On the other hand, the tear strength values were 66.3 and 26.3 kN/m for the CNT and N-990 CB containing compounds. Almost three times increase in the tear strength yielded only a negligible improvement in the abrasion resistance. The same authors [15] reported successful improvement in the abrasion resistance of a medium crosslinked HNBR elastomer when both MWCNT and CB were at the
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same time incorporated in the matrix (50 phr N-550 type CB). When only nanotubes were added to the HNBR mixture, without the presence of the CB, the tear strength of the nanocomposite was more increased. On the other hand, the abrasion resistance was less when the HNBR mixture contained only nanotubes without CB. Pal et al. [43] prepared and tested natural rubber (NR)/styrene–butadiene rubber (SBR) (with high styrene contents) compounds). The mixes contained 40 phr CB (type of the CB was varied: N-231 and N-774) both with 10 phr ENR epoxidized natural rubber/Cloisite 20A (ratio 1:1) compounds. DIN (ASTM 5963-04), Du-Pont (ASTM D394) abrasion tests were performed on standard devices and on a home built machine, where the rotating rubber wheels (mounted with aluminum core) were pressed against different rock materials (sandstone, concrete and granite) with constant normal loads. The authors observed that for the standard DIN and Du-Pont wear tests the OC containing rubber mixtures showed better wear resistance than the traditional compounds without OC fillers. Against different rock materials the weight loss trend of the investigated rubbers differed from that received in standard abrasion tests. Against sandstone and concrete the organoclay containing rubbers were more wear resistant than those with traditional fillers. On the other hand, against granite, all the rubber compounds exhibited similar wear resistances.
2.1.2 Thermoplastic Rubbers A polyolefin elastomer, namely poly(ethylene–octene) copolymer (Engage 8401; referred further as TPO thermoplastic elastomer), was filled with OC and MWCNT, respectively. The coding, nanofiller content and the mechanical properties of the related nanocomposites is disclosed in Sect. 2.2.2. Abrasive pin(rubber)-on-plate (abrasive) type (POP) tests (cf. Fig. 2) were performed to test the abrasive wear behavior of the TPO-based nanocomposites containing OC and MWCNT fillers. The 1 mm thick elastomer composite specimens were glued to the surface of a fixed ceramic ball (Ø * 8.77 mm). The so prepared specimens were pressed against the plate counterbody which was covered with sandpaper (P 180, average grind size 82 lm). The radius of the circular path was 15 mm, the sliding speed was 20 cm/s, the normal load was 1 N and the sliding distance was 100 m. During the tests the COF was online registered and the loss volumes of the different materials were compared with each other based on the related diameter of the wear paths (cf. Fig. 2). The worn surface of the specimens was inspected by optical microscopy. The wear behavior of the TPObased nanocomposites is collated in Fig. 3. Considering the diameter of the wear track the OC containing nanocomposites exhibited increased wear resistance in contrast to the MWCNT filled ones the wear of which was enhanced. With increasing OC content the wear resistivity of the composites decreased. So the lowest wear was observed for the TPO-2 (containing 2 wt% OC) composite. Among the MWCNT-containing materials, the lowest wear was measured for the
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Fig. 3 Diameter distribution of the wear tracks after abrasive POP tests for the TPO based nanocomposites containing OC and MWCNT in different amounts. Designations: TPO-2, -5, -10 contain 2, 5 and 10 wt% OC whereas TPO-C05, -C1, -C2 contain 0.5, 1 and 2 wt% MWCNT, respectively
TPO-C1 (containing 1 wt% MWCNT). These results suggest that the nanofiller content may have an optimum value, which should be determined accordingly. The time dependence of the coefficient of friction (COF) values during the abrasive POP tests shows a fast increase initially, and a pretty stable value in the steady state (cf. Fig. 4). The fast COF increase may be caused by the temperature rise at the contact surface. The OC nanofiller does not influence the COF value, whereas with varying MWCNT content the COF varies, as well. The smallest COF value, and at the same time the smallest wear for the TPO/MWCNT series, were measured for the TPO-C1 (containing 1 wt% MWCNT). Compared to this nanocomposite, the steady state COF value and the wear increased for both TPO-C05 and TPO-C2 composites. The abrasive wear patterns of the TPO-based nanocomposites are very similar to each other. Figure 5 displays a light microscopic picture which was taken from the TPO-5 (containing 5 wt% OC) nanocomposite. Here a typical unidirectional abraded pattern can be resolved. Compared to the worn surface of conventional thermoset rubbers in this thermoplastic rubber rather plastic deformation dominates instead of coarse edged cracking. The surface is covered by grooves and large size wear particles (cf. Figs. 5, 6). During sliding large amount of heat is generated and accumulated by the friction supporting the plastic deformation of the material under adiabatic conditions. This may cause the melting of the TPO, reflected by enhanced irreversible deformations. Accordingly, due to the thermoplastic behavior of TPO its worn surface differs from that in Fig. 1, which is more characteristic for conventional thermoset rubbers.
2.2 Sliding-Type Sliding wear gained importance nowadays, because efforts are made to avoid the lubrication between the sliding rubber and other elements. This trend is reasoned by economical and environmental issues. Therefore, investigations on the sliding type wear under dry conditions are very important. Recall that for sealing and
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Fig. 4 Coefficient of friction (COF) as a function of sliding time for the TPO/OC (left side) and TPO/MWCNT (right side) nanocomposites during the abrasive POP tribotest. For designations cf. Fig. 3
related applications the elastic and viscoelastic properties of the rubbers are exploited. In these applications usually high adhesion arises between the rubber and the counterbody owing to their smooth surfaces. Under relative displacements between the contacting pairs adhesive sliding and wear takes place. To meet the industrial needs new rubber formulations, novel coatings and rubber–polymer sliding pairs have to be developed. For sliding wear tests the ball-on-prism configuration is often favored [60, 61], as in the related device allows to test many specimens at the same time. In this ball(counterbody)-on-prism(rubber) test usually a revolving steel ball is pressed against two rubber plates forming a prism. Disadvantage of this technique is that no friction force can be detected. Other test possibility is the roller-on-plate (ROP; also termed to shaft-on-plate (SOP) test) configuration, where a rotating shaft is
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pressed against a rubber plate, whereby the normal and friction forces are detected (see Fig. 7). Generally it is not always easy to distinguish between the different methods being differently classified in respect to the wear type. The simplified laboratory POP test (see Fig. 7) can be labeled as adhesive type sliding tribotest if
Fig. 5 Light microscopic picture of the worn surface of the TPO-5 (containing 5 wt% OC) nanocomposite. Note: sliding direction is downward
Fig. 6 Wear mechanism concluded for the thermoplastic polyolefin elastomer containing OC and MWCNT nanofillers
Fig. 7 Schematic set-up of the tribotesting devices used. Designations: pin(steel)-on-plate(rubber) (POP), roller(steel)-on-plate(rubber) (ROP)
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the surface of the abrader counterbody is smooth enough. However, there is no straightforward definition on how to differentiate between abrasive-like and ‘‘smooth’’ surfaces, and thus between abrasive and sliding wears.
2.2.1 Conventional Rubbers The sliding, friction and wear behaviors of peroxide cured hydrogenated nitrile rubber (HNNBR) with 10 and 30 phr MWCNT were investigated [62]. Mechanical properties (hardness, tensile modulus, ultimate tensile strength and strain, tear strength) of the rubbers were determined. Dynamic mechanical thermal analysis (DMTA) was also performed and the apparent crosslink density estimated. The related mechanical and material related data are summarized in Table 1. Friction and wear characteristics were determined in pin(steel)-on-plate(rubber) (POP) configuration in which a steel pin (100Cr6; arithmetical roughness, Ra, less than 1 lm) with a hemispherical tip of 10 mm diameter rotated along a circular path, see Fig. 7. The pin was pushed against the rubber plate with a given load. The following parameters were selected for this configuration––normal load: 2 N, sliding speed: 250 mm/s, duration: 1.5 h (h). Measuring both the normal and the friction force components via a torque load cell the COF values were calculated and monitored during the test. To study the sliding wear a further test, viz. ROP, was also used, see Fig. 7. A rotating steel roller (100Cr6, diameter: 10 mm, width: 20 mm, Ra & 0.9 lm) pressed against a rubber strip of 9 mm width in the related tribotester. The frictional force induced by the torque was measured online and thus the COF was registered during the test. The test parameters were––load: 2 N; sliding speed: 250 mm/s; duration: max. 1.5 h. It is noteworthy that the selected test parameters were adjusted to the praxis by considering the contact pressure and sliding speed.
Table 1 Network-related and mechanical data for the HNBR rubbers containing MWCNT as filler. For the description of the test procedure see reference [62] HNBR HNBR ? 10phr MWCNT HNBR ? 30phr MWCNT Mc [g/mol] nc [10+26 9 m-3] tan d at Tg [-] Density [g/cm3] Shore A [] Martens hardness [MPa] M-0.01 DMTA [MPa] M-100 [MPa] M-200 [MPa] Tensile strength [MPa] Tensile strain [%] Tear strength [kN/m]
1931 3.3 1.36 1.057 42 1.45 3.9 1.27 1.3 4.4 280 4.2
1113 5.7 1.1 1.057 56 2.16 6.7 2.4 2.58 4.9 179 8.0
150 45.3 0.5 1.135 82 10.2 52.4 12.7 12.4 16.5 132 22.0
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Figure 8 shows the effect of MWCNT on the COF. One can recognize that the COFs are lower in POP tests, while they are markedly higher for the ROP configuration. This is due to the related test set-up favoring substantial heating under ROP. For the latter case even rubber against rubber type friction may be observed (when the roller is coated with rubber). Moreover, the change in the COF with the MWCNT content is different when various test devices are used. For the POP configuration the COF is increasing with increasing MWCNT content compared to the neat HNBR. Using the ROP device the measured COF is smaller when MWCNT is added to the HNBR, but with increasing MWCNT content the COF is increasing. Comparing the specific wear rate results one can see, that their differences may cover some order of magnitudes (cf. Fig. 8). Higher wear is caused by the POP than by the ROP tests. Using MWCNT additives the wear performance of the related HNBR is always better. With increasing filler content the resistance to wear of the HNBR compound is enhanced. It should be born in mind that the specific wear rate is usually decreasing with increasing active filler content as demonstrated for CB-containing ethylene–propylene–diene (EPDM) rubbers [63, 64]. Characteristic scanning electron microscopic (SEM) pictures taken of the worn tracks of the HNBR stocks are given in Figs. 9, 10 for the POP and ROP tests, respectively.
Fig. 8 Measured specific wear rate (Ws). a and steadystate coefficient of friction (COF) values and b for the POP and ROP tribotests
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Fig. 9 SEM pictures taken from the wear tracks after POP tests. a pure HNBR, b HNBR ? 10phr MWCNT, and c HNBR ? 30phr MWCNT. Note: sliding direction is downwards [62]. (Reprinted from: Felhös, D. Karger-Kocsis, J. Xu, D.: Tribological testing of peroxide cured HNBR with different MWCNT and silica contents under dry sliding and rolling conditions against steel. Journal of Applied Polymer Science, 2008, 108, pp. 2840–2851, Copyright 2010, with permission from John Wiley and Sons)
POP tests: The surface of the evaluated wear track on the pure HNBR shows that quite large particles were chipped off from the surface and a crater-like pattern appeared (Fig. 9a). At 10 phr MWCNT content, however, the so-called Schallamach waviness [65, 66] becomes recognizable (Fig. 9b). With further increase of the MWCNT content the Schallamach waves become less resolved and a band-like pattern appeared (Fig. 9c). The reason of the latter may be linked with the MWCNT fragments induced reinforcing effect. Nevertheless, the smooth worn surface suggests low wear, that was found in fact (cf. Figs. 8a, 9c). Interestingly, the smooth surface appearance was not accompanied with low COF (cf. Fig. 8b). ROP tests: On the worn surface of the pure HNBR chipped off debris, roll formation and spherical particles are seen (Fig. 9a). This suggests high COFs (cf. Fig. 8b) and high wear rates (cf. Fig. 8a), which were found, in fact. By reinforcing the HNBR with 10 phr MWCNT the worn surface become smeared (Fig. 10b) indicating for lower COF and specific wear rate (Ws) in accordance with the experimental findings. With increasing MWCNT content (30 phr) the surface of the wear track become ever smoother, and practically no wear debris or cracks
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Fig. 10 SEM pictures taken from the wear tracks after ROP tests. a pure HNBR, b HNBR ? 10phr MWCNT, and c HNBR ? 30phr MWCNT. Note: sliding direction is downwards [62]. (Reprinted from: Felhös, D. Karger-Kocsis, J. Xu, D.: Tribological testing of peroxide cured HNBR with different MWCNT and silica contents under dry sliding and rolling conditions against steel. Journal of Applied Polymer Science, 2008, 108, pp. 2840–2851, Copyright 2010, with permission from John Wiley and Sons)
could be found (Fig. 10c). In that case the MWCNT worked probably as a solid lubricant. Recall that values of both the COF and specific wear rate strongly depend on the testing rigs. This is in line with the usual claim that friction and wear are ‘‘system properties’’. Increasing MWCNT content generally enhances the resistance to wear. This finding is similar what was reported on the effect of CB fillers in rubbers. Between the mechanical and tribological properties of the investigated materials no definite correlations can be found. The complexity of the wear behavior requires a large data base in respect with the thermo-mechanical and network-related properties of rubbers based on which their effects on wear may be traced. Xu and Karger-Kocsis published a well founded study on the tribological properties of organophilic layered nanosilicate (OLS) filled HNBR [67]. As clay modifier a quaternary amine with hydroxile and double bonds was used. The properties of the OLS filled nanocomposites were compared to the pure
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HNBR material. Dry sliding tests were performed applying different testing configurations. It was found, that for all the executed dry sliding tests, the wear resistance was negatively influenced by the intercalated OLS filler. On the other hand, the coefficient of friction was slightly increased or dramatically decreased depending on the testing configurations. Gatos et al. [68] performed an extensive wear study on montmorillonite (MMT with and without organophilic treatment) filled HNBR and EPDM rubber grafted with maleic anhydride (EPDM–MA) elastomers. The MMT modifiers were the following: octadecyl amine (ODA; primary amine), octadecyl trimethylammonium (ODTMA) and methyl-tallow-bis(2-hydroxy-ethyl) quaternary ammonium salts (MTH). It was found that for dry sliding, all the MMT additives resulted in an increased coefficient of friction. The measured wear was higher for the HNBR/ MMT nanocomposites than for the neat HNBR rubber, except the unmodified MMT containing HNBR. For the EPDM–MA with MMT–ODA and MMT– ODTMA both the coefficient of friction and the measured specific wear were reduced under POP condition.
2.2.2 Thermoplastic Rubbers Two thermoplastic rubber compounds were investigated in respect to their dry sliding friction and wear characteristics. One of them is a polyester-based linear thermoplastic polyurethane (Elastollan C60A W; referred further as TPU), which according to the datasheet contains 10 wt% plasticizer, has a nominal Shore A hardness of 64, and the melting temperature of hard segments is at about 185C. The second one was a polyolefin-based version, viz. TPO. These two thermoplastic rubbers were filled with unmodified MWCNT (NC7000 of Nanocyl, Sambreville, Belgium) and organoclay (OC, in this case: Nanofil 5, Süd-Chemie AG, München, Germany). The surfactant of the latter was distearyl dimethyl ammonium chloride. The OC containing TPU and TPO thermoplastic elastomers are denoted further as TPU-2, -5, -10 and TPO-2, -5, -10, where the last number means the wt% of the OC filler. The MWCNT containing TPUs and TPOs are designated as TPU-C05, -C1, -C2 and TPO-C05, -C1, -C2, where C stays for CNT and the last number denotes the wt% of the MWCNT filler. The composites were prepared on an open mill at 145C (TPO) and at 185C (TPU). 1 mm and 4 mm thick sheets were hot pressed at 150C (TPO) and at 190C (TPU). Standard ISO527-2 specimens were cut from the 1 mm thick sheets for tensile tests. The Shore A hardness was tested on the 4 mm thick samples. The mechanical properties of the thermoplastic elastomers and their nanocomposites are summarized in Table 2. In Table 2a the mechanical properties of the TPO based nanocomposites are listed. The Shore A hardness of the TPO increased with incorporation of both the OC and MWCNT additives. The tensile strength of the TPO nanocomposites decreased with increasing MWCNT and OC. The decrease in the tensile strength is larger for MWCNT than for OC. Interestingly the change in the strain to failure value is minor for the OC additive. In contrary to this,
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Table 2 Mechanical properties of the thermoplastic elastomers and their nanocomposites Shore-A Tensile strength [MPa] Strain to failure [%] Density [g/cm3] a TPO-0 TPO-2 TPO-5 TPO-10 TPO-C05 TPO-C1 TPO-C2
85 86 88 89 87 88 89
9.61 8.61 8.69 8.78 7.89 6.68 8.35
829 790 823 820 702 554 777
0.884 0.890 0.902 0.931 0.884 0.891 0.893
TPU-0 TPU-2 TPU-5 TPU-10 TPU-C05 TPU-C1 TPU-C2
74 74 75 74 75 77 79
29.85 28.63 22.98 13.56 21.57 17.70 22.83
1003 967 897 739 817 651 864
1.163 1.168 1.176 1.192 1.164 1.165 1.171
b
the strain to failure diminishes markedly with the incorporation of MWCNT. The density of the nanocomposites increases monotonously with increasing additive content. In Table 2b the mechanical properties of the TPU-based nanocomposites are collated. Note that the OC filler slight increases the Shore A hardness. In contrast, the MWCNT filler is a more ‘‘active’’ one, which increases the Shore A hardness with its increasing content. The tensile strength and the strain to failure values both decrease with incorporation of OC and MWCNT. Decreasing tendency in the strength and elongation properties is to observe with increasing OC and MWCNT contents. The density of the nanocomposites increases monotonously with increasing additive content. Sliding friction tests were made on a POP-type tribometer (see Fig. 7) (THT-0000 type from CSM, Peseux, Switzerland). For the sliding the 4 mm thick samples were used. The TPU-based nanocomposites were tested with ceramic (Al2O3) ball of 8.77 mm diameter, applying 15 N normal load (FN) and a sliding speed of 25 cm/s with a distance of 400 m. The radius of the circular path was 4.99 mm. The TPO-based nanocomposites were tested using Cr6 bearing ball of 6 mm diameter applying 5 N load 10 cm/s and 200 m. The radius of the circular path was 4.99 mm. The loss volume of the specimens after POP tests was estimated based on the half with and on the depth of the wear grooves cross section, which was considered to be a half ellipse. To get the loss volume the estimated area of the half ellipse was multiplied with the circumference of the orbital revolution of the pin (r = 4.99 mm). Loss volume (wear) results after POP tests are shown for the TPO/OC and TPO/ MWCNT nanocomposites in Fig. 11. Though the results underlay a large scatter, it can be established, that both OC and MWCNT additives improve the wear resistance. For the TPO/OC and TPO/MWCNT nanocomposites the lowest wear
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Fig. 11 Loss volume (wear) results in POP tests for the TPO-based nanocomposites containing OC and MWCNT
was measured at 2 and 1 wt% additive, respectively. This suggests again that enhanced sliding wear resistance is probably linked with an optimum amount of nanofiller. Considering the development of COF as a function of time (cf. Fig. 12) some preliminary conclusions can be drawn on the wear process. At the very beginning of the tests there is a large drop in the COF values. Recall that this is often the case for POP tests. The increasing heat and the damage of the specimen surface along with the appearance of wear debris are the main reasons for the COF decrease in the initiation phase (t \ 100 s). In the following time period (steady state) the COF changes marginally. The second drop in the COF versus time curve may represent the total destruction of the material structure at the surface. At this stage large sized cracks appear at the surface and the final wear groove pattern develops. One exception is the TPO-C2 composite (see diagram (b) in Fig. 12), the COF of which is steeply reduced shortly after the initial phase. After the total surface destruction, the cracks propagate fast at the surface and the COF values tend to level off at ca. 0.5. The length of the steady state correlates with the wear resistance of the material, as the related time intervals are in accord with the loss volume results. The longest steady state period (as well as the lowest wear) among the TPO/OC and TPO/MWCNT nanocomposites were found at the 2 wt% OCand 1 wt% MWCNT-containing samples (TPO-2 and TPO-C1, respectively). Similar observations were reported in Ref. [69] where the COF and wear rate of the investigated thermoplastic elastomers were decreased with increasing hardness of the material. In [30] this observation was limited for the POP configurations. In Fig. 13 one can see the worn surface of the TPO based nanocomposites after POP tests. These photographs display cracks across the wear path which are opened in the sliding direction. It can be established that the dominant wear mechanism does not change, but the density of the cracks increases with the addition of nanofillers. Based on the optical inspection the following wear mechanism occurs in POP test (see Fig. 14). First, Schallamach type waves form at the surface due to the high adhesion and low modulus of the elastomer (cf. Fig. 14a). At the wave fronts, associated with large deformations, cracks initiate (cf. Fig. 14b). The initial cracks
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Fig. 12 Coefficient of friction (COF) during POP tests for the TPO-based nanocomposites containing OC and MWCNT
Fig. 13 Light microscopic pictures taken of the worn surfaces of TPO/OC and TPO/MWCNT nanocomposites after POP tests. a plain TPO, b TPO-5, and c TPO C2. Note: sliding direction is downward
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at the surface are plastically deformed owing to ‘‘over-sliding’’ of the rigid counterbody (cf. Fig. 14c). During the repeated ‘‘over-sliding’’ the cracks open in the sliding direction and the cracks undergo severe plastic deformation. This process results in the formation ‘‘rubber tongues’’ which is typical for sliding wear patterns in rubbers (cf. Fig. 14d). In Fig. 15 the loss volume results for the TPU/OC and TPU/MWCNT nanocomposites are collected. One can see that the incorporation of the nanofillers reduced the wear resistance of the composites. The incorporation of MWCNT
Fig. 14 Wear mechanism of the TPO nanocomposites for POP test. a formation of Schallamach waves, b initiation of cracks due to the Schallamach waviness, c deformation and growth of the initial cracks, d formation of rubber tongues at the surface
Fig. 15 Loss volume (wear) results after POP tests for TPU/OC and TPU/MWCNT nanocomposites
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seems to be less ‘‘harmful’’ than OC from the view point of the wear. Interestingly the same OC and MWCNT amounts (2 and 1 wt%, respectively) yield the most wear resistant composites similar to the TPO-based ones. On the other hand, all of the nanofiller containing composites are less wear resistant than the neat TPU matrix. The course of COF as a function of sliding time is depicted in Fig. 16 for the TPU based nanocomposites. Again, a large drop in the COF value can be observed in the initial phase of the sliding friction tests. After this fast reduction the COF increases with time (and thus sliding distance). The smallest COF during the tests in the steady state was observed for the neat TPU matrix material. The highest steady state COF values were observed for the TPU-5 (cf. red curve in diagram (a) in Fig. 16)
Fig. 16 COF as a function of sliding time during POP test for TPU/OC (a) and TPU/MWCNT nanocomposites (b)
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and for the TPU-C05 (cf. blue curve in diagram (b) in Fig. 16). Nanocomposites with the highest COFs exhibited the largest loss volumes, as well. For the TPU based composites, particularly for the MWCNT filled ones, the COF becomes rather stable for a given period after the initial wear-in phase. The length of the stable period correlates again with the wear resistance. The lowest wear among the MWCNT filled nanocomposites was measured for the TPU-C1 and this sample possessed also the longest time interval with stable COF. This may imply that the onset of surface cracking occurred after more wear cycles in this material than in the others. Figure 17 shows the light microscopic pictures taken of the worn tracks of the TPU (cf. Fig. 17a, b) and TPU/OC (cf. Fig. 17c, d) and TPU/MWCNT nanocomposites after POP tests (cf. Fig. 17e, f). Similar wear mechanisms can be recognized on these figures for all of the composites. One of these wear formations is the Schallamach-type waviness which is characteristic for elastomeric materials, especially at low hardness (cf. Fig. 14). The other group of wear formations belongs to the plastic deformations. These are namely the ‘‘ploughing’’ and the extrusion- or ‘‘ironing’’-type deformations, which are characteristic for thermoplastic materials (cf. Fig. 18a, b). The plastic deformations are partly generated by surface roughness peaks of the counterbody, partly due to the high adhesion between counterbody and rubber substrate. Interestingly, multiplied Schallamach waviness appears on the surface after the POP tests. This means, that between the large sized Schallamach-type waves smaller ones are also present. Plastic deformations are also present in the wear tracks, namely ploughing, extrusion and ‘‘ironing’’ (cf. Fig. 19). The Schallamach waviness and the clearly recognizable plastic deformations suggest that the investigated elastomer has a dual (rubber- and thermoplastic-like) feature at the same time.
Fig. 17 Light microscopic pictures taken of the worn surfaces of TPU and TPU/OC and TPU/ MWCNT nanocomposites after POP test. Designations: a, b plain TPU-0 matrix; c, d TPU-5, and e, f TPU-C2
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Fig. 18 Plastic micro-deformations on the worn surface schematically. Designations: a ploughing and b ironing, extrusion
Fig. 19 Combined wear mechanisms (dual waviness and plastic deformation) in the TPU based nanocomposites after POP test
2.3 Rolling-Type Pioneering works in the rolling friction for rubbery materials have to be credited to Tabor, Greenwood and coworkers in 1950s [70–76]. They all attributed the rolling friction principally to mechanical hysteresis in the substrate when the latter is deformed and the deformation is released by the passage of the rolling body [77]. When a spherical indenter is pressed into rubber a certain amount of elastic work is performed. As the indenter moves forward elastic work is needed to deform the rubber in front of the indenter whilst it is recovered from the rear. Assuming ideal elastic behavior, the rubber behind the indenter would yield identical amount of work and no net kinetic energy would be lost. But in fact, a constant fraction of the input elastic energy is lost as a result of viscoelastic hysteresis in the rubber. During rolling this is the primary source of the frictional work [71]. In addition, the adhesion phenomenon is superposed to the hysteretic losses. In special cases both effects can be comparable with each other. In Fig. 20 the adhesion
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Fig. 20 The adhesion effect in case of a rolling cylinder on elastomeric plane [79] (a). Wear mechanism (b) and wear formation (c) in case of rolling [80] tribotests due to repeated stresses (fatigue)
phenomenon between a rolling cylinder and an elastomeric plane under unlubricated conditions is depicted. The distance j between the rigid body and the elastomeric plane represents the molecular roughness of the elastomer. One can assume, that the polymer chains are fixed in the bulk at their one end (via crosslinking) while their other is free. During sliding the free molecule ends touch the surface of the rigid body and adhere to it. The detaching process needs more energy than the adhering process. As a consequence, the molecules at the detaching (peeling) site of the contact region will be extended causing additional friction forces [78]. Based on this adhesion the same process causes friction even if the rigid body in Fig. 20 is a surface asperity in micron- or in nano-range in sliding movement [79]. Different wear mechanisms can be observed, when the wear of rubbery materials is tested under rolling. In case of lubricated rolling contact, the typical wear process is of fatigue type. During fatigue-like tests, if the surfaces are lubricated, peculiar subsurface crack initiation and propagation may take place. It is known, that in Hertzian-like contacts the maximum stresses and strains appear in the bulk near to the surface. So, the probability of crack initiation is the highest there, where the arising strains reach their maximum. Thus cracks arise below the surface and the cracks propagate during the fatigue process. This crack propagation advances as long the cracks grow together and wear particles break out of the surface (see Fig. 20b). This phenomenon is called pitting or gouging. In Fig. 20c one can see the schematic picture of the gouging effect, the onset of which can be observed in fretting or in rolling tests. The rolling friction and wear properties of elastomeric materials can be determined in different testing configurations, such as the orbital rolling ball-onplate (O-RBOP) test. For this test, however, one needs to clarify the kinematics of the rolling ball, which qualitatively determinates the wear and friction during the tests. In Fig. 21 one can see the schematic sketch of the O-RBOP test configuration
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Fig. 21 Schematic sketch of the orbital rolling ball(steel)on-plate(rubber) test configuration (O-RBOP)
Fig. 22 Superposition of the sliding velocities in the contact area during orbital rolling [62]. (Reprinted from: Felhös, D. Karger-Kocsis, J. Xu, D.: Tribological testing of peroxide cured HNBR with different MWCNT and silica contents under dry sliding and rolling conditions against steel. Journal of Applied Polymer Science, 2008, 108, pp. 2840–2851, Copyright 2010, with permission from John Wiley and Sons)
wherein a steel ball, which is pressed with a constant normal load into the rubber specimen, rolls along an orbital path on the rubber specimen (cf. Fig. 21). The investigation of the kinematical process of the rolling tribotests is inevitable, due to the fact, that under dry rolling conditions, beneath the fatigue process, additional sliding wear occurs via microslipping. This significant microslipping (sliding of the non-adhering counter surfaces) can take place because of the differences in the elastic moduli of the rolling pairs, and due to superposition of different rolling phenomena. Accordingly, the dry rolling is always accompanied with sliding wear. Note that depending on the surface quality of the counterbody the sliding wear may turn into abrasive one, similar to the wear of car tires. In Fig. 22 one can see the superposition of different sliding processes during O-RBOP test in the wear track. The different sliding processes are originated from the different components (drilling, forward rolling) of the orbital rolling method. The detailed description of the rolling–sliding processes during the O-RBOP tests is given in Ref. [62].
2.3.1 Conventional Rubbers Different rubber blends based on HNBR and FKM with and without CB/MWCNT fillers were tested with the O-RBOP tribotest configuration under dry rolling
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conditions. The peroxide curable HNBR was mixed with 20 phr CB (N550), 10/ 20/30 phr MWCNT (Baytubes C 150 P from Bayer MaterialScience, Leverkusen, Germany), respectively. The MWCNT filler in the composites will be designated with the letter ‘‘C’’. The samples prepared are referred to as HNBR, HNBR-20CB, HNBR-10/20/30C, respectively. The digits in the designations represent the contents of fillers in phr. HNBR with hybrid fillers (i.e. CB and MWCNT) was also produced. The corresponding composition contained 20 phr CB and 10 or 20 phr MWCNT, and denoted as HNBR-CB-C. Being both peroxide-curable HNBR and FKM (Viton GF-600S of DuPont Performance Elastomers Geneva, Switzerland) were combined in the following HNBR/FKM ratios: 100/0, 100/33.3 and 100/100 (designations: HF-1:0, HF-1:1/3 and HF-1:1 respectively). To the above mixtures 10 phr MWCNT was incorporated on a laboratory mill. Note that the cure recipe was not adjusted, i.e. the ready-to-cure HNBR was just diluted by FKM. This means that a direct comparison can only be made between rubber compounds with and without MWCNT filler when their HNBR/FKM ratio is the same. The rubbers involved are further on referred to as HNBR, HNBR-10C, HF-1:1/3, HF-1:1/3-10C, HF-1:1, and HF-1: 1-10C, respectively. Their network-related parameters, density and hardness values are summarized in Table 3. From the point of the physical properties it is an important question how the CNTs are dispersed in the matrix. To investigate the distribution of the nanotubes, transmission electron microscopic (TEM) pictures were taken from the prepared HNBR/FKM/MWCNT composites. The TEM inspection showed that the MWCNT filler is preferentially located in the continuous HNBR phase (see
Fig. 23 TEM picture of the HNBR/FKM/MWCNT (HF1:1-10C) nanocomposite. Note that MWCNT is located only in the HNBR phase
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Fig. 23). Attention should be paid to the fact that the HNBR is the continuous phase even in the HF-1: 1mixture. In the home-made O-RBOP test configuration (cf. Fig. 21), the rubber sheet is worn by a steel ball (100Cr6, diameter: 14 mm, arithmetical roughness Ra: 1 lm), rolling along a circular path diameter: 33 mm), pushed by a defined normal load against the rubber sheet. The normal load (FN) was 90 N, the revolution number is 280 rpm, and the test duration was 3 h. The specific wear rate was computed according to Eq. 1. The loss volume (DV) (cf. Eq. 1) was calculated by measuring the depth and width of the wear track and estimating it as a half ellipse. The cross section of the wear track was assessed by a white light profilometer. The wear degree can be assessed by specific wear rate (Ws): Ws ¼
DV mm3 FN L N m
ð1Þ
where DV [mm3] is the volume loss, FN [N] is the normal load, L [m] is the overall rolling distance. Table 4 summarizes the COF and Ws of HNBR with and without fillers measured in O-RBOP tests. The COF changes marginally with the increasing content of MWCNT (cf. Table 4). Using MWCNT the wear resistance of the related HNBR mixes was strongly enhanced, compared to that of the pure HNBR. The results suggest that incorporation of CB, lowered the Ws and slightly enhanced the COF. CB proved to be better in improving the rolling wear resistance of HNBR compared to MWCNT. The Ws values of HNBR–FKM hybrids with and without MWCNT measured in O-RBOP show a great improvement in the wear resistance. Most prominent is the enhancement of the wear resistance for the HNBR/FKM 100-33.3 MWCNT hybrid (HF-1:1/3-10C; here the Ws is 30 times lower than that is for the neat HNBR). Incorporation of MWCNT into HNBR/FKM decreased the Ws and increased the COF. The COF values of the HNBR/FKM hybrids are higher compared to the neat HNBR. The SEM photos in Fig. 24 were taken from the wear track of the HNBR, after O-RBOP tests. For the pure HNBR rubber Schallamach type pattern with roll head can be found on the both outer and inner sides of the wear track (cf. Fig. 24). The wave fronts in the two regions are adverse to each other. This reflects that the direction of the ball movement in the two regions is opposite (cf. Fig. 22). The waves are especially well-developed in the outer side, confirming a larger speed in the outer compared to the inner region (cf. Fig. 22). In the middle section large agglomerates can be observed (cf. Fig. 24). Note that the effect of the spin of the ball is negligible in the centre region. The debris produced in the outer and inner regions were likely swept into the centre region by the spin of the ball favoring the formation of such agglomerates. Based on the SEM pictures one can draw the following conclusions: for CB and MWCNT, the Schallamach-type pattern completely disappeared and the debris fragmented surface in the centre region of the wear track result in a lower Ws and higher COF compared to the pure HNBR, respectively (see Ref. [62]).
Mc [g/mol] mc [mol/m3] tan d at Tg Density [g/ cm3] Martens hardness [MPa]
1176 899 1.11 1.057
407 2710 0.68 1.102
161 7052 0.50 1.135
1775 615 1.25 1.091
407 2710 0.68 1.102
3088 371 1.16/0.31 1.146
1068 1213 0.36/0.56 1.296
3715 352 0.56/0.64 1.309
HF-1:1
1182 1119 0.35/0.58 1.322
HF-1:1-10C
1.02 ± 0.56 2.20 ± 0.24 3.37 ± 0.71 10.23 ± 0.84 0.66 ± 0.29 3.37 ± 0.71 1.19 ± 0.09 2.21 ± 1.08 1.12 ± 0.06 2.98 ± 0.75
2013 525 1.37 1.057
Table 3 Network-related properties and hardness of the HNBR/FKM–CB/MWCNT nanocomposites and hybrids HNBR HNBR-10C HNBR-20C HNBR-30C HNBRHNBRHF-1:1/3 HF-1:1/320CB 20CB-20C 10C
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HNBR10C
HNBR20C
HNBR30C
2.6 ± 0.739 1.05 ± 0.122 0.917± Ws [10-4 mm3/ 2.95 ± 0.839 Nm] COF[-] 0.042 ± 0.01 0.038 ± 0.00181 0.043 ± 0.0015 0.041±
HNBR 1.05 ± 0.122
HNBR20CB-20C
HF-1:1/3-10C
HF-1:1
0.283 ± 0.0512 0.066 ± 0.0279 0.788 ± 0.430
HF-1:1/3
0.117 ± 0.0154
HF-1:1-10C
0.0497 ± 0.0004 0.043 ± 0.0015 0.049 ± 0.0004 0.053 ± 0.0004 0.055 ± 0.0016 0.055 ± 0.0026
0.613 ± 0.201
HNBR20CB
Table 4 Wear and coefficient of friction (COF) results for the HNBR/FKM–CB/MWCNT nanocomposites and hybrids
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Fig. 24 SEM photos taken from the rolling wear track after Orbital-RBOP test of the HNBRPure. Picture to the left is the outer region, the picture in the middle is the centre region and the picture to the right was taken from the inner region. Note: rolling direction is downward the arrows indicate the local sliding direction [62]. (Reprinted from: Felhös, D. Karger-Kocsis, J. Xu, D.: Tribological testing of peroxide cured HNBR with different MWCNT and silica contents under dry sliding and rolling conditions against steel. Journal of Applied Polymer Science, 2008, 108, pp. 2840–2851, Copyright 2010, with permission from John Wiley and Sons)
It is noticeable that the wear loss of the filler reinforced HNBR system correlates with the occurrence of the Schallamach-type pattern. The dramatically prohibited waviness corresponds to the decreased Ws for CB and MWCNT reinforcement. The introduction of active fillers (MWCNT and CB) is the reason for the less developed or lacking waves, as often observed in rubbery systems with active fillers (e.g. [64, 81, 82]). One may also observe that incorporation of 10 phr MWCNT enhances only slightly the wear resistance of the HNBR. However, when 20 phr fillers were added, the specific wear rate was strongly reduced. Interestingly, only very slight further decrease of Ws was found when the compounds containing higher amounts of filler. SEM photos taken from the rolling wear tracks of HF-1:1/3 with and without MWCNT after Orbital-RBOP tests are shown in Fig. 25. In the outer region of HF-1:1/3, Schallamach type pattern appears with some fragments (cf. Fig. 25a). Clusters from debris (agglomerates) were formed in the centre and inner regions (cf. Fig. 25c). Incorporation of MWCNT changes the basic wear mechanisms in the outer region (cf. Fig. 25b). Fatigued-induced formation of holes becomes the main wear mechanism instead of the Schallamach wavy pattern in HF-1:1/3. Hole development and debris clustering were found in the centre region of HF-1:1/3 10C (cf. Fig. 25b). Fatigue-induced pittings and small ‘‘ironed’’ particles are seen in the inner region of the HF-1:1/3 10C (cf. Fig. 25d). The wear mechanisms for the HF-1:1 (MWCNT) did not change practically when compared to HF-1:1/3 (MWCNT)––cf. Fig. 26. However, the inner region of HF-1:1 is full with ‘‘ironed’’ rolls, debris (cf. Fig. 26c). By comparing the worn surfaces of HNBR–FKM with and without MWCNT, one can notice that the Schallamach wavy pattern disappears after MWCNT introduction. This change in the wear mechanism reflects a reduction in the specific wear rate, which was found also for other active fillers, e.g. CB. It is worth noting that the Schallamach-type wear disappears with increasing reinforcement of a rubber independent of the type of the active filler used [67, 83].
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Fig. 25 SEM photos taken from the rolling wear tracks of HF-1:1/3 without (a, c) and with (b, d) MWCNT (HF-1:1/3-10C) after Orbital-RBOP tests. Designation: a outer region, b outer and centre regions, c centre and inner regions, and d inner region. Note: rolling direction is downward [84]. (Reprinted from: D. Xu, J. Karger-Kocsis, Z. Major and R. Thomann: Unlubricated rolling wear of HNBR/FKM/MWCNT compounds against steel. Journal of Applied Polymer Science, 112, (2009), pp. 1461–1470, Copyright 2010, with permission from John Wiley and Sons)
2.3.2 Thermoplastic Rubbers The TPU rubbers used were as follows: a polyetherdiol-based TPU (Estane 58300, Shore A = 80) with and without 5 wt% organoclay (Nanomer I30P), and a polyesterdiol-based TPU (Estane 58142, Shore D = 65) with and without 5 wt% Nanomer I30P. The surfactant (organophilic modifier) of Nanomer I30P was octadecyl amine. Nanocomposites were produced by incorporating the OC in the TPUs via extrusion melt compounding. The four materials used are further on referred to as TPU–L, TPU–L/OC, TPU–H and TPU–H/OC in the text. Note that L and H stand for the low and high hardness TPU grades, respectively. The wear topography in the outer region of TPU–L/OC 150 N sample is very unique (cf. Fig. 27b). It is of Schallamach type with rolls, but some ‘‘waves’’ are locally aligned parallel to the rolling direction. So, Fig. 27b shows the onset of a two-dimensional wavy network, mesh elements of which are oriented along the
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Fig. 26 SEM photos taken from the rolling wear tracks of HF-1:1 without (a, c) and with (b, d) MWCNT (HF-1:1-10C) after Orbital-RBOP tests. Designation: (a, b) outer and centre regions, (c, d) inner region. Note: rolling direction is downward [84]. (Reprinted from: D. Xu, J. Karger-Kocsis, Z. Major and R. Thomann: Unlubricated rolling wear of HNBR/FKM/MWCNT compounds against steel. Journal of Applied Polymer Science, 112, (2009), pp. 1461–1470, Copyright 2010, with permission from John Wiley and Sons)
rolling direction. This peculiar Schallamach-type pattern may probably explain the lower specific wear rate of this OC loaded sample compared to that of the neat TPU–L (cf. Fig. 27a and b; Table 5). SEM pictures in Fig. 27a support why higher specific wear rate and COF were found for the plane TPU–L sample. Schallamach pattern with rolls (‘‘tongues’’) are well discernible in the outer and inner regions. Table 5 Specific wear results for the TPU–L/OC and TPU–H/OC nanocomposites after O-RBOP tests Normal load/test duration Specific wear rate (mm3/N m) 60 N/3 h 90 N/3 h 120 N/3 h 150 N/3 h
TPU–L
TPU–L/OC
TPU–H
TPU–H/OC
1.01E–04 1.12E–04 1.20E–04 1.69E–04
1.45E–05 3.94E–05 6.53E–05 1.06E–04
6.26E–06 1.39E–05 1.67E–05 2.88E–05
5.89E–06 4.71E–06 1.49E–05 2.66E–05
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Fig. 27 SEM pictures taken from the rolling wear track of TPU–L (a) and TPU–L/OC (b) after O-RBOP test performed at 150 N normal load for 3 h [85]. (Reprinted from: D. Xu, J. KargerKocsis and A. K. Schlarb: Rolling friction and wear of organoclay-modified thermoplastic polyurethane rubbers against steel. Kautschuk, Gummi, Kunststoffe, 61 (2008), pp. 98–106, Copyright 2010, with permission from Hüthig GmbH)
In the first approximation, one can state that the better developed the Schallamachwaves are, the higher the COF (cf. Table 6) and specific wear rate are (cf. Table 5). In the outer region of TPU–H an embryonic wavy pattern appears after dry rolling test (cf. Fig. 28a). This is believed to be the primary stage of the Schallamach pattern formation. No track edges and wavy pattern could be resolved in the worn surface of TPU-/OC (cf. Fig. 28b). On the other hand, the wear track becomes different from the unfilled version when OC is incorporated. The Schallamach pattern is the more pronounced the lower the hardness of the related rubbers is. Note that this claim holds also for traditional rubbers ([64] and references therein). The OC modification improved the resistance to dry rolling wear (cf. Table 5). The improvement was the higher the lower the hardness of the parent TPU was. The COF was slightly reduced by OC incorporation when tested under the same conditions for softer TPUs. However, the COF was increased in presence of OC for the harder TPUs (cf. Table 6). It is noteworthy that no such clear tendencies could be concluded for the dry sliding wear of OC-modified rubbers [68]. The COF was hardly affected by the test duration. By contrast the specific wear rate was reduced with the time. This can be explained by a reduction Table 6 Coefficient of friction (COF) results for the TPU–L/OC and TPU–H/OC nanocomposites Normal load/test duration COF (-) 60 N/3 h 90 N/3 h 120 N/3 h 150 N/3 h
TPU–L
TPU–L/OC
TPU–H
TPU–H/OC
0.0366 0.0347 0.0375 0.0438
0.0288 0.0315 0.0352 0.0383
0.0172 0.0240 0.0259 0.0341
0.0193 0.0265 0.0306 0.0371
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Fig. 28 SEM pictures taken from the rolling wear track of TPU–H (a) and TPU–H/OC (b) after O-RBOP test performed at 150 N normal load for 3 h [85]. (Reprinted from: D. Xu, J. KargerKocsis and A. K. Schlarb: Rolling friction and wear of organoclay-modified thermoplastic polyurethane rubbers against steel. Kautschuk, Gummi, Kunststoffe, 61 (2008), pp. 98–106, Copyright 2010, with permission from Hüthig GmbH)
in the real contact surface owing to debris formation. Both COF and specific wear rate increased with increasing normal load (cf. Tables 5, 6). Recall that the wear mechanisms in Orbital-RBOP tests are complex due to the superimposed rolling and spinning.
3 Outlook and Future Trends Novel nanofillers, such as organophilic-modified clays (OC) and carbon nanotubes (CNT) are promising additives to improve the tribological performance, targeting both reduced coefficient of friction and wear rate, of rubbers. Further work is needed, however, to realize this potential. The results reported above suggest that OC and CNT are less suited for rubber/rubber and rubber/thermoplastic blends because of their selective embedding into one component of the compositions. CNT, when not properly functionalized and thus dispersed, may be far less effective than expected. In this respect attention should be paid to the fact that the quantitative assessment of the nanofillers’ dispersion is not solved. Hybrid reinforcing of rubbers, i.e. the combined use of traditional and novel fillers, may be the right research strategy for the next future. Less information is available on how these nanofillers affect the lubricated wear of rubbers. Very interesting is the tribological (both dry and lubricated) performance of nanofilled rubbers under high-pressure conditions. According to the authors’ opinion both OC and CNT may be suitable additives to improve the resistance to rapid gas decompression of rubbers. Further impetus to the friction and wear of rubbers is expected from modeling studies considering the wear types of rubbers accordingly. Though considerable efforts are devoted to this issue, an in depth understanding is still
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lacking. Note that model predictions would be very helpful for further material (recipe) development and ‘‘refining’’ of tribological tests. Acknowledgments Dr. Dávid Felhös is very thankful to Mr. György Szabó and his family their selfless and friendly advocacy to take a fresh start in Miskolc. The authors express their thanks to Dr. Dan Xu for the performed tests and for the results presented in Sects. 2.3.1 and 2.3.2 and to Dr. Kálmán Marossy for the preparation of the TPU and TPO based nanocomposites.
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Index
A abrasion, 4–6, 13, 58, 180, 201–203, 205–208, 213, 218, 226–231, 320, 322, 329, 343, 347–349, 379 abrasion resistance, 4, 58, 180, 202, 205–208, 213, 218, 224, 226, 322, 324, 329, 347–349 actuators, 281, 283, 287, 293, 304, 305 additive, 58, 61, 106, 113, 155, 157, 161, 162, 164, 165, 167–175, 186, 203, 208, 209, 215, 223, 270, 278, 294, 310, 354, 357–359, 375 aggregation, 23, 57, 58, 65–67, 70, 71, 73, 81, 97, 142, 220, 293–296, 304 alkylammonium, 180 aspect ratio, 6–10, 12, 14, 29, 48, 52, 53, 89, 99, 120, 140, 144, 145, 151, 220, 234, 242, 243, 247, 253–256, 260, 262, 264, 309, 322, 327, 328, 332, 344, 346
B barrier effect, 49, 158, 162, 163, 170, 173 barrier function, 180 barrier performance, 23, 234, 236, 248 basal plane spacing, 236, 240–242, 247, 248 blending, 11, 96, 103, 104, 116, 156, 163, 208, 209, 213, 227, 228, 331, 332, 344 broadband dielectric spectroscopy, 89, 94, 114, 116
C capacitance, 33–35, 113, 285, 286, 290, 300 carbohydrates, 136
carbon black, 5–8, 13, 29, 38, 39, 50, 52, 54, 55, 58, 69, 83, 90, 91, 113–116, 120, 136, 139, 140, 141, 153, 179, 180–184, 189, 190, 192, 193, 195, 198, 201, 203, 206–208, 210, 214–217, 220–223, 226–228, 260, 293, 294, 302, 304, 344, 379 Carbon nanofibres, 158, 163, 317 Carbon nanotubes, 5, 7, 8, 12–17, 23, 50, 53, 54, 90, 96, 97, 115, 117, 120, 130, 131, 136, 158, 163, 164, 172, 175–177, 184, 187, 188, 193–195, 198, 210, 267, 271, 281, 310, 317, 332, 340, 343, 344, 375–378 catalytic effect, 158, 167, 171 chain movement, 135 chemical architecture, 238, 252 chemical modifications, 235 cloisite, 166, 201, 215, 238, 242, 271, 274, 311, 349 coefficient of friction, 205, 267, 343, 350, 351, 354, 357, 360, 370, 374, 375 compatibilizer, 104, 117, 166–168, 180, 185–187, 196, 213 compression molding, 217, 218, 333 Coulombic interaction, 121, 124 crack growth resistance, 180 cross-linked, 32, 61–65, 71, 75, 76, 333, 341 crystallisation kinetics, 310 curing, 49, 70, 72, 77, 79, 91, 115, 140, 201, 248, 315, 329, 332, 333, 335, 345, 348, 376
381
382 D damping treatment, 308 degradation, 13, 48–50, 54, 65, 156–159, 161–163, 165, 170, 172, 177, 179, 180, 185–188, 190, 191, 195–197, 206, 217, 221, 223, 226, 233, 247, 250, 251, 255, 265–268, 271, 274, 275, 277, 278, 280, 317, 318, 329, 330, 338, 339, 342 degree of bonding, 92 dielectric, 17, 31, 33–35, 54, 67, 89, 94, 95, 99, 101, 102, 106, 108–114, 116–118, 190, 281, 284–288, 290, 291, 303–305, 377 dielectric elastomers, 284, 286, 287, 291, 305 dielectric loss peaks, 95 DIN abrasion, 218, 224, 225, 248 disordered solids, 91 dispersion, 3, 5, 9–19, 21, 23, 26, 48, 52–54, 58, 61, 68–71, 73, 75, 77, 78, 92, 98, 102, 103, 117, 118, 130, 136, 142, 146, 155, 156, 158, 159, 162, 165, 167–171, 174–177, 180, 189, 193, 195, 211–213, 217, 218, 222, 226, 240, 256, 259–265, 270, 271, 294, 295, 322, 328, 330, 344, 375, 376 dispersion state, 155, 159, 165, 168, 174, 175, 211 Du-Pont abrasion, 218, 224 durability, 4, 185, 308, 339 dynamic mechanical, 5, 17, 31, 37, 38, 40, 43, 51, 54, 55, 68, 71–74, 77, 80, 89, 90, 93, 95, 96, 102, 104, 112, 114, 115, 255, 267, 274, 278, 315, 332, 345, 353, 377 dynamic mechanical analysis, 31, 68, 71, 73, 77, 80, 89, 90, 95, 102, 104, 112, 114, 115, 315, 332
E ecological tendencies, 282 elastomeric nanocomposites as biomaterials, 264 elastomeric nanocomposites for biomedical applications, 259 electrical behaviour, 89, 109, 110, 131 electrical circuitry, 286 electrical devices, 155 electrical field, 89, 282, 287–289, 291, 296–304 electro active polymers, 281, 282
Index electrospun, 212, 311, 312, 328, 339 energy harvesters, 281, 303 ENR, 49, 70, 71, 74, 76, 90, 103–105, 176, 201, 213–215, 217, 219, 220, 222, 224, 226, 346, 349, 379 entropy modulus, 135 epoxy, 16, 52–54, 70, 74, 78, 79, 84, 114, 176, 177, 233, 234, 239–252, 254–256, 276, 277, 280, 325, 310, 341, 376–379 equilibrium state, 37, 90 exfoliation, 10, 12, 23, 110, 131, 165, 236, 241–243, 247, 248, 253, 255, 256, 262, 276, 313, 320, 321, 344, 376 extrusion, 120, 309, 313, 363, 364, 372
F fibres, 82, 158, 163, 197, 212, 260, 307–314, 317, 325, 327, 331, 333, 336, 337, 340 filler shape factor, 58 finite element approach, 254 fire hazards, 156 fire regulations, 155 flame inhibition, 157 flame retardancy, 155–157, 159, 161, 162, 164, 169, 170, 172–178, 180, 189, 328, 347, 378 flame retardant, 155–163, 165, 167–178 flammability, 155, 156, 159–162, 170, 171, 174–177, 190, 197, 256, 317 fractography, 29 fracture properties, 308, 328 frequency, 32, 33, 36, 44–46, 54, 93–96, 98–101, 103, 106–111, 172, 290, 291
G gas permeability, 136, 255, 276, 330 green rubber, 59, 60 grinding, 4, 11, 14
H harvesting, 281, 283, 284, 286–288, 291, 292, 301–305 hybrid, 14, 43, 44, 51, 52, 57, 59, 62, 71, 74, 77–85, 103, 116, 117, 131, 189, 190, 207, 234, 276, 279, 318, 365–368, 373, 374, 376 hydrolytic sol-gel process, 59, 81 hydrophilicity, 77, 243, 272
Index I Impedance, 3, 32–34, 54, 94, 271, 290, 302 in situ, 10, 55, 57, 58, 60–66, 68–85, 103, 116, 131, 138, 154, 162, 188, 190, 213, 261, 267, 272, 320, 330, 331, 342, 344, 346, 376 incompatibility, 58, 209, 233, 234, 238, 242, 247, 318 injection molding, 120 inorganic oxides, 57, 58, 60, 80, 81 interatomic, 119, 120, 121, 127 intercalation, 10, 32, 51, 99, 111, 131, 166, 167, 177, 213, 219, 226, 233, 235, 242, 248, 260–262, 272, 276, 318, 320, 322, 330, 341, 346, 378 interface, 12, 13, 16, 43, 50, 53, 54, 59, 71, 78, 82, 91, 101–103, 105, 106, 108–110, 113–116, 119, 131, 146, 148, 191, 201, 206, 209, 233, 234, 238, 250, 264, 271, 275, 295, 308, 312, 327, 338 interface area, 119 intumescence, 157
L layered silicate, 5, 11, 49, 51, 52, 63, 90, 98, 99, 114, 116, 117, 131, 157, 158, 162, 163, 167, 175, 185, 196, 197, 213, 255, 256, 260, 271, 276, 277, 313, 318–320, 322, 328, 339–342, 346, 378 Lennard-Jones function, 124 lifecycle, 179 loss modulus, 38–40, 45, 47, 93 loss tangent, 34, 37–39, 42, 44, 97, 98, 103
M magnetic force microscopy (MFM), 23 mechanical milling, 137, 138 Mechanical performance, 97, 103, 109, 191, 233–235, 238, 308, 331 melt flow, 120 miniaturelization, 280 modification, 15, 61, 75, 84, 98, 108, 110, 131, 136, 146, 155, 157, 161, 162, 167, 171, 174, 190, 196, 233–235, 237–254, 256, 272, 333, 335, 374, 376 mold geometry, 120 molecular and continuum modeling, 119 molecular chain entanglement, 73, 122, 142, 221, 318
383 molecular dynamic effects, 89 montmorillonite, 9, 18, 51, 52, 90, 98, 115–117, 120, 136, 151, 157, 162–165, 167, 172–178, 180, 185, 190, 195–197, 214, 233, 235, 236, 238, 239, 241–245, 247–253, 255–257, 260, 275, 276, 279, 311, 325, 339–342, 357, 378 Moore’s Law, 283 multi-walled carbon nanotube, 96, 115, 117, 135, 141, 142, 158, 163, 193, 194, 317, 376, 377
N nanoclay, 9–11, 19, 99, 115–117, 136, 154, 164, 177, 180, 181, 183, 185–187, 189, 191, 213–215, 217, 219, 222, 224, 310, 313, 317, 318, 325, 326, 378 nanocomposite characterization, 262 nanographite, 3, 7, 10, 11, 17–21, 26, 27, 29, 30, 32–36 nanostructured materials, 307 nanotechnology, 91, 131–133, 198, 259, 278, 307, 338, 376 natural rubber, 10, 15, 41, 51–53, 55, 66, 70, 82, 83, 90, 95, 114–118, 135–137, 139, 149, 152–154, 160, 176, 189–192, 201, 206–208, 213–215, 219, 220, 223, 227, 228, 346, 349, 376–378 Nyquist plots, 33, 35
O oligomer, 59, 74, 90, 103, 331, 332, 342 opacity, 64 optimisation, 309 organoclay, 51, 52, 98, 99, 102, 104, 105, 115–117, 154, 177, 190, 201, 212, 213, 226, 255, 256, 276–278, 311, 312, 314, 320–322, 339–341, 343, 346, 347, 357, 372, 374, 375, 378, 379 orientation, 12, 18, 24, 30, 41, 50, 54, 106, 108, 113, 128, 132, 135, 136, 145, 146, 148, 151, 153, 210, 211, 235, 262, 307, 309, 322, 345, 346
P packaging, 155, 190, 192, 197, 233–236, 238, 239, 243, 247, 259, 329 particulate releases, 179
384
P (cont.) percolation, 3, 35, 37, 39, 54, 63, 291, 344, 345 phase changes, 89, 94, 112, 113 piezo devices, 281 piezoelectric materials, 282, 283 poly(aryl-ether–ether–ketone), 307, 334 polydimethylsiloxane, 160, 186, 197, 258, 271 poly(p-phenylene benzbisoxazole), 337 polyaniline, 307, 329, 330, 342 polyarylacetylene, 332, 333, 342 Polymer Nanocomposite Synthesis, 261 polyurethane, 9, 49, 51, 80, 85, 90, 103, 114, 116, 117, 131, 146, 160, 161, 164, 176, 187, 189, 196, 197, 233–240, 242, 243, 254, 255, 257, 258, 260, 265, 269, 276–280, 287, 294, 302, 304, 305, 307, 318, 325, 326, 340–342, 357, 374, 375, 377–379 polyurethane nanocomposites, 235–239, 242, 243, 254, 255, 269, 278, 280, 304, 305, 340–342, 377 potential energy, 119–121, 123, 124, 127, 133, 283 potential energy functions, 119, 121, 123 precursors, 8, 60, 65, 74, 81 processing conditions, 234, 309 processing technique, 58, 335 protective barrier, 158, 159 proteins, 136, 137 pseudo bilayer, 250 pultrusion, 120 pyrolysis, 50, 157, 182, 187, 188, 191, 192, 197, 333
R radius of gyration, 10, 12 recycling, 179–189, 191, 192, 195–197 relaxation phenomena, 89, 91, 93–95, 98, 99, 102, 106, 110, 111, 113, 114 reliability, 308 renewable, 281, 283 resource cascading, 179, 181, 183–186, 188, 191 retreading, 182, 183, 212 rolling wear, 368, 371–375, 379 Rubbing, 14, 207
S scanning electron microscopy, 24, 217, 218 scavenging, 281, 283
Index Schallamach wave, 355, 361 service life, 205, 308, 338 shape memory alloys, 282, 283 sliding wear, 205, 350, 351, 353, 359, 361, 366, 374, 378, 379 sol-gel, 53, 57, 59–62, 65, 66, 70, 71, 73–75, 77–85, 189, 277, 293, 313, 336, 342 solid state physical properties, 119 solvent intercalation, 10, 32, 51, 99, 111, 131, 166, 167, 177, 213, 219, 226, 233, 235, 242, 248, 260–262, 272, 276, 318, 320, 322, 330, 341, 346, 378 specific wear rate, 343, 354–356, 368, 371, 373–375 steric effect, 158 stiffness, 5, 11, 12, 33, 58, 67, 121–123, 180, 223, 256, 258, 265, 309, 310, 313, 314, 316, 320, 324, 326, 328, 330, 335–337, 344–346 storage modulus, 16, 38–47, 71–73, 77, 80, 93, 97–99, 103–106, 267, 321 strain-induced crystallization behavior, 153 stress-strain curves, 36, 67, 69, 73, 74, 76, 138, 141, 142, 144, 147, 149, 151, 327, 335 structure-property relationship, 71, 82, 83, 89, 340 styrene-butadiene rubber, 51, 73, 84, 115, 116, 131, 135, 139, 170, 206, 349, 376 substrate, 62, 98, 115, 158, 164, 169, 170, 207, 235, 247, 249, 317, 363, 364, 379 surfactant, 16, 17, 54, 73, 98, 136, 165, 167, 174, 260, 278, 357, 372, 378 suspensions, 16, 91, 241, 247, 248 swelling, 9, 17, 49, 51, 60, 61, 65, 67, 71–76, 215, 216, 260, 275, 321, 339, 345 synergies, 155, 171, 175
T tear strength, 5, 58, 79, 180, 182, 220, 221, 223, 346, 348, 349, 353 tensile mode, 93, 98 tensile tests, 217, 264, 269, 332, 335, 336, 357 thermal degradation, 48, 65, 156, 158, 159, 163, 197, 223, 250, 251, 255, 275, 317, 342 thermal resistance, 156, 160, 164, 217 thermally stable, 155, 156, 161, 168, 174, 330
Index thermoforming, 120 thermomechanical, 60, 89, 255, 276, 277, 338 titania, 8, 58, 60, 64–66, 80–82, 180, 194 torsional angle, 121, 123 TPE gels, 3, 17, 26, 28, 29, 43, 45–49 Transmission Electron Microscopy (TEM), 165, 217 transitions, 39, 93, 282 twin-screw extruder, 312, 313, 335 tyres, 179–184, 186, 187, 189, 191, 192, 198, 201–208, 212, 213
V van der Waals interactions, 25, 344 vermiculite, 233, 245, 247–249, 256 viscoelastic, 31, 32, 43, 55, 93, 96, 132, 309, 321, 322, 341, 346, 347, 351, 364, 378, 379 vulcanization time, 58
385 W water vapor permeation, 236–238, 242, 243, 249 Wear, 49, 96, 109, 114, 115, 186, 188, 189, 201–208, 213, 214, 226–229, 283, 291, 329, 336, 337, 342, 343, 347–359, 361–366, 368–373, 378, 379 wide-angle X-ray diffraction, 137
X X-ray Diffraction (XRD), 167, 262, 317 X-ray scattering, 57, 60, 64, 65, 153 XRD, 18, 80, 90, 99, 110, 167, 216, 219, 220, 222, 226, 257, 262, 263, 266, 274, 317, 330
Z zirconia, 58, 60, 64, 66, 81, 82
