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7Mechanical and Viscoelastic

Characterization of MultiphasePolymer Systems

Poornima Vijayan P.

School of Chemical Sciences, Mahatma Gandhi University, Kerala, India

Siby Varghese

Rubber Research Institute of India, Kottayam, Kerala, India

Sabu Thomas

Centre for Nanoscience and Nanotechnology, Mahatma Gandhi University, Kerala, India

7.1 Introduction

Multiphase polymer systems, which form a fast-growing development in the area of polymer science andtechnology, include blends, composites, alloys, interpenetrating polymer networks (IPNs) and gels. Polymerblends are mixtures of two or more polymers and/or copolymers in which the minor component contributesat least 2 wt%. Polymer blends have gained significant growth in the last two decades and they constitute ca.36 wt% of total polymer consumption. Current worldwide market volume for polymer blends and alloys isestimated to be more than 700,000 metric tons/year with an average growth rate of 6 to 7 % [1–6]. Figure 7.1illustrates the price/performance of commercial polymer blends.

Commercial polymer blends have existed for decades; however, the concepts of miscibility, phase behaviorand the basic nature of polymer blends were not well understood. It is perhaps that the successful com-mercialization of miscible blends of poly (2,6-dimethyl-1,4 phenylene ether) and polystyrene prompted theinterest in polymer blends. The inspiration for polymer scientists and industrialists to focus on polymer blendsrather than synthesizing new materials arose from cheaper development costs, maximum diversification and

Handbook of Multiphase Polymer Systems, First Edition. Edited by Abderrahim Boudenne, Laurent Ibos, Yves Candau, and Sabu Thomas.© 2011 John Wiley & Sons, Ltd. Published 2011 by John Wiley & Sons, Ltd.



252 Handbook of Multiphase Polymer Systems

PERFORMANCEPRICE

CONSUMPTION

SPECIALITY POLYMERS

PEEK, LCP, PES, PSU, PAR, PEI, PPS

PA, PBS, PET, PC, POM, PPO, ABS, SMA, Acrylics

HDPE, LDPE, LLDPE, PP, PS, HIPS, PVC

ENGINEERING THERMOPLASTICS

COMMODITY THERMOPLASTICS

Figure 7.1 The price/quality profile of commercial polymer blends. Reprinted from [1]. Copyright (2002) withpermission from Springer.

increased use of existing polymers. Additional benefits are: (1) improved specific properties; (2) improvedmeans for the industrial recycling; (3) plant flexibility and productivity; and (4) diversified applications ofpolymer blends which include automotive, electronics and electrical, building and constructions, medical andpackaging, etc.

Recently, nanostructured polymer blend systems have become increasingly important. In nanoblends, thescale of dispersion of one polymer into another is in general below 100 nm.Nanostructured polymer blendsvery often exhibit unique properties that are directly attributed to the presence of structural entities havingdimension in the nanometer range [7].

On the other hand, interpenetrating polymer networks (IPNs) are a blend of two or more polymers in anetwork form, at least one of which is synthesized and/or crosslinked in the immediate presence of the other(s).When one of the polymers is crosslinked, the product is called semi IPNs. An IPN can be distinguished frompolymer blends, blocks, or grafts in two ways: (1) an IPN swells, but does not dissolve in solvents, and (2) creepand flow are suppressed [1]. The IPN is held together exclusively by mutually entangled bands of nanometerscale dimension. This kind of ‘supramolecular’structure favors shear yielding, which is the major energydissipation mechanism in thermosets [8]. Existence of the physical interlocking (interlocked macrocycles ofthe two component networks) had been the major factor in determining the characteristic properties [9]. IPNstructuring offers unique possibilities for composites with respect to fiber-to-matrix adhesion [8].

Another important and widely-used multiphase polymer system is polymer gels. A polymer gel consistsof an elastic crosslinked network and a fluid filling the interstitial spaces of the network. The network of longpolymer molecules holds the liquid in place and so gives the gel what solidity it has [10]. Polymer gels canbe easily deformed by external stimuli, and generate force or execute work on the external environment. Ifsuch responses can be translated from the microscopic level to a macroscopic scale, a conversion of chemicalfree energy into mechanical work should be realized. The ability of polymer gels to undergo substantialswelling and collapsing as a function of their environment is one of the most remarkable properties of thesematerials. The phenomenon of gel volume transitions, which can be induced by temperature, pH, or ionicstrength, has prompted researchers to investigate gels as potential actuators, sensors controllable membranesfor separations, and modulators for delivery of drugs [11].



Mechanical and Viscoelastic Characterization of Multiphase Polymer Systems 253

Polymer science uses the concept of a composite as a material formed by combining different phases. Inconventional polymer composites, many types of inorganic fillers with dimensions in the micrometer range,e.g., calcium carbonate, glass beads, etc. have been used extensively to enhance the mechanical properties ofpolymers. A nanocomposite is defined as a composite material where at least one of the dimensions of one ofits constituents is in the nanometer scale. How nanoparticles affect the properties of nanocomposite materialsis of concern to engineers and scientists who are interested in the applications of nanocomposites [12].

Final properties of the multiphase polymer system are strongly controlled by various factors such asmorphology, dispersion, mixing conditions, filler content, etc. This chapter discusses the various parametersthat affect the mechanical properties of the above-mentioned multiphase polymer systems separately. Variousmodels have been used to predict the mechanical and viscoelastic properties.

7.2 Polymer Blends

7.2.1 Ultimate Mechanical Properties and Modeling

Major parameters which control the mechanical properties of the polymer blends are: morphology, mixingconditions, blend composition, crosslinking, compatibilization and filler addition, etc.

7.2.1.1 Influence of Blend Morphology on Mechanical Properties

The final properties of polymer blends are greatly dependent on the morphology of the systems. Morphologyof a polymer blend indicates the size, shape and spatial distribution of the component phases with respect toeach other. Since it is well established that most of the properties of the polymer blend are strongly influencedby the type and fineness of the phase structure, the study of the control of the morphology of the polymerblend has emerged as an area of continuous interest to polymer material scientists in the last few decades[13, 14].

For a given blend, various types of useful morphologies (Figure 7.2) for different end properties such ashigh strength and toughness, toughness coupled with stiffness, good barrier properties, and high flow can beobtained [14]. From the point of view of broader classification, multiphase polymer blends may be dividedinto two major categories:

1. Blends with a discrete phase structure (i.e. droplets in matrix).2. Blends with a bicontinuous phase structure (i.e. co-continuous).

Blends with a discrete phase structure are most common in which droplets of the minor phase dispersed in amatrix of the major phase. These types of blends are often used in rubber modification of brittle polymers. Theminor phase can also be dispersed as fibers, for example in self-reinforcing polymer blends. In these kinds ofblends, the properties are mainly improved in the direction of the fibers. In co-continuous morphologies, theinteresting feature is that both components, in all directions, can fully contribute to the properties of the blend.

Attempts have been made by Thomas and co-workers to correlate the blend morphology with the observedmechanical properties [15, 16]. Varghese et al. studied the correlation between morphology and mechanicalproperties of nitrile rubber/ethylene-vinyl acetate copolymer (NBR/EVA) [15]. Morphology of the NBR/EVAblends indicated a two-phase structure in which the minor phase is dispersed as domains in the majorcontinuous phase. However, between 40 and 50 wt% of NBR content both NBR and EVA exist as continuousphases and generate a co-continuous morphology. The stress-strain curves of EVA-rich blends have a similar




Drops Double emulsion Laminar

(barrier)(toughness and stiffness)(toughness, surfacemodification)

Ordered microphasesCocontinuousFibers

(high flow, electricalconductivity,toughness, stiffness)

(strength, thermal expansion)

1 μm

10 μm

Figure 7.2 Schematic of useful morphologies of polymer blends. Reprinted from [8]. Copyright (2006) Taylorand Francis Group.

behavior as that of pure EVA, NBR rich blends as that of typical unfilled rubber and the 50–50 blend with aco-continuous morphology showed an intermediate stress–strain behavior. Tensile strength and elongation atbreak increase with the increase in EVA content. The increase was found to be sharper when the EVA contentwas more than 40% where it formed a continuous phase.

Kumar et al. found that the morphology of the blend has a strong influence on the mechanical properties ofnylon/NBR blend [16]. The mechanical properties were found to increase rapidly beyond 40 wt % of nylon(Figure 7.3). This abrupt rise in mechanical properties is associated with the fully co-continuous nature ofthe nylon matrix (Figure 7.4).

Veenstra et al. measured and compared the mechanical properties of polymer blends with co-continuousphase morphologies to the properties of blends of the same polymers with a droplet-matrix morphology[17]. In their study, PS/poly(ether-ester) and polypropylene/styrene-butylene-styrene (PP/SEBS) copolymerblends were prepared with both morphologies (dispersed-matrix and co-continuous). Elastic moduli of the co-continuous blends were significantly higher than the moduli of the dispersed blends. However, no significantdifference in tensile strength was found when the co-continuous blends were compared to blends withdroplet-matrix morphology.




50

40

30

MA

XIM

UM

TE

NS

ILE

ST

RE

NG

TH

, MP

a20

10

00 20 40 60

WEIGHT % OF NYLON

80 100

Figure 7.3 Effect of nylon/NBR blend composition on maximum tensile strength. Reprinted from [16]. Copyright(1996) with permission from John Wiley & Sons.

7.1 μm

7.7 μm

20.3 μm

(a) (b)

(c)

Figure 7.4 SEM images of the blend morphology of (a) N30 showing the dispersed nylon particles in thecontinuous NBR matrix. (b) N50 showing a co-continuous morphology. (c) N70 here NBR particles are dispersedin the continuous nylon matrix. Reprinted from [17]. Copyright (2000) with permission from Elsevier.




2000

1000

100

10

1

Ey

(MP

a)

0.0

(b)

0.1 0.2 0.3 0.4

Volume fraction PP

0.5 0.6 0.7 0.8 0.9 1.0

3000

1000

Ey

(MP

a)

100

300.0 0.1 0.2 0.3 0.4

Volume fraction PS

0.5 0.6 0.7 0.8 0.9

(a)

1.0

Figure 7.5 (a) Young’s moduli for the PS/poly(ether–ester) systems Ia (�) and Ib (•) as a function of volumefraction PS; (b) Young’s moduli for the PP/SEBS systems IIa (�) and IIb (•) as a function of volume fraction PP(when the error bars are not visible they are smaller than the marker) [17].

In Figure 7.5a, the Young’s moduli for two blend systems of PS/poly(ether-ester), Ia (the blends processedat 230◦C) and Ib (the blends processed at 200◦C), are plotted as a function of PS. Morphology of the Iashowed co-continuity over a small composition range (50 to 60 vol % PS), and the morphology of the Ibblends showed a broad range of co-continuity (30–60vol %PS) (Figure 7.6). All other compositions of Ia andIb blends showed droplet-matrix morphologies. Moduli show a high increase when PS becomes continuousthroughout the sample. It is important to note that when the moduli of the co-continuous blends are comparedto that of the dispersed morphologies (with the same volume fractions), it becomes evident that at low volumefractions of PS, the co-continuous blends show higher values for the Young’s modulus than the dispersedblends. It is very clear that PS contributes more to the modulus of the blend when it is continuous than when itis dispersed in the poly(ether-ester) matrix. This is associated with the fact that in co-continuous morphologiesboth phases take part in the load-bearing process. At higher volume fractions of PS, the difference in modulusrelated to the morphology diminishes. Similarly, in Figure 7.5b, the Young’s moduli for two blend systemsof PP/SEBS blends, IIa (the blends processed at 250◦C), and IIb (the blends processed at 190◦C), are plottedas a function of the volume fraction of PP. The morphology of the IIa blends showed continuity over asmall composition range (50 to 60 vol% PP), and the morphology of the IIb blends showed a broad rangeof co-continuity (40 to 80 vol % PP). From this, it is evident that co-continuous blends show a much highervalue for Young’s modulus than the dispersed blends.

7.2.1.2 Influence of Mixing Conditions on Mechanical Properties

The mechanical properties of polymer blends strongly depend on processing conditions (temperature, time,intensity and type of mixing and nature of flow). Mechanical properties of natural rubber/poly(methylmethacrylate) blends were investigated as a function of mixing conditions by Oommen and Thomas [18].A comparison of tensile strength and tear strength of the solution-cast and melt-mixed samples was carriedout. Even though both systems exhibit considerable improvement in mechanical strength, the melt-mixed




Figure 7.6 SEM micrographs of PS/poly(ether–ester) blends of systems Ia (a and b) and Ib (c and d) with 30and 40 vol.% PS [17].

samples were found to show lower strength values as compared to respective solution-cast samples (Table 7.1).Melt- mixed blends exhibit lower strength both in the compatibilized and uncompatibilized states comparedto the solution-cast system. The improved strength of the solution-cast blend is because molecular-levelmixing is achieved during solution mixing and this leads to improved adhesion between the phases. Thehigh shearing action (80 rpm) and temperature (145◦C) employed during the preparation of the blend by themelt-mixing technique might have caused degradation of NR and PMMA, resulting in substantial reduction instrength.

7.2.1.3 Blend Composition

Blend composition of the multiphase polymer system is one of the important factors which controls themorphology of the system and hence the mechanical properties. The mechanical properties of PLA–PCLblends can be tuned through the blend composition. Blends of biodegradable polymers poly (d,l-lactic acid)and poly(e-caprolactone) were prepared by Broz et al. [19]. Figure 7.7 shows the strain-at-failure data acrossthe entire mass fraction range of PLA–PCL blend. This value decreased monotonically as the mass fractionof PLA increases. A rather precipitous drop was seen from 0 to 0.6 followed by flat behavior thereafter. Thisindicates that even small amounts of glassy PLA are capable of embrittling the PCL matrix; at PLA mass




Table 7.1 Mechanical properties of 50/50 NR/PMMA blends with graft copolymer [18]. Standard deviationvalues are given in parentheses.

Solution-Cast System Melt-Mixed System

GraftCopolymer(%)

TearStrength(N/mm)

Elongationat Break

(%)

TensileStrength(N/mm2)

TearStrength(N/mm)

Elongationat Break

(%)

TensileStrength(N/mm2)

0 9.80 322 3.75 9.46 49 1.25(1.62) (33.6) (1.06) (0.58) (3.52) (0.1)

5 20.21 185 9.61 14.75 31 2.01(1.64) (13.6) (0.74) (0.51) (3.16) (0.12)

10 22.40 80 12.37 17.76 21 3.10(1.35) (3.28) (0.63) (0.32) (2) (0.34)

15 23.47 132 13.37 − 24 3.14(1.26) (4.83) (1.16) (1.4) (0.19)

fraction 0.2 the strain-at-failure has decreased 50% relative to pure PCL. This may be due to blending ofthe glassy PLA into the PCL matrix or simply to the formation of PLA inclusions in the blend with someinterfacial adhesion. Flat part of the curve at PLA mass fractions at and above 0.6 was consistent with theformation of a continuous PLA matrix; PLA is glassy and the strain-at-failure is expected to be insensitiveto strain because it depends on molecular parameters such as the free volume in the matrix phase. If thishypothesis were true, it suggests that there is very little mixing between PLA and PCL in this compositionrange; otherwise, the PCL would have been expected to have a toughening effect on the blend. Therefore, itappears that there may be some association of PLA and PCL at low PLA mass fraction but at higher PLAcontents there is little or no reinforcement due to blending.

0.40

Str

ain-

at-f

ailu

re

Mass fraction PLA

0.35

0.30

0.25

0.20

0.15

0.10

0.05

0.000.0 0.2 0.4 0.6 0.8 1.0

Figure 7.7 Plot of strain-at-failure across the composition range for the PLA/PCL blends. Reprinted from [19].Copyright (2003) with permission from Elsevier.




40

Yie

ld S

tres

s (M

Pa)

0.0

Mass fraction PLA

0.2 0.4 0.6 0.8 1.0

35

30

25

20

15

10

Figure 7.8 Plot of yield stress as a function of composition in PLA–PCL blend. Reprinted from [19]. Copyright(2003) with permission from Elsevier.

In Figure 7.8, the yield stress data as a function of composition are shown. The yield stress is insensitiveto composition from PLA mass fraction 0–0.4 then increased linearly up to 1.0. This linear dependence ofyield stress at compositions ranging from pure PLA down to PLA mass fraction 0.6 suggests PCL blendingin this regime simply dilutes the PLA matrix like the presence of voids in the material, reducing the totalstress necessary to fracture the samples. The flat trend at compositions ranging from pure PCL to PLA massfraction 0.4 suggests there must be some reinforcement due to interactions between PLA and PCL, otherwisea similar reduction in yield stress in blended samples would be observed as in the PLA-rich samples.

Above a threshold PLA mass fraction of 0.4, the modulus and ultimate tensile strength increased almostlinearly as a function of composition. This threshold may be due to strengthening of the blend interface inthis regime.

George et al. studied the effect of the blend composition on mechanical properties of SBR/NR blends [20].Mechanical properties of homopolymers and blends are given in Table 7.2. The SBR/NR blends are denotedby N0, N30, N50, N70, and N100, where the subscripts denote the weight percent of NR in them. The propertiessuch as tensile strength and elongation at break were increased from N0 to N100. Mechanical strength ofSBR increased upon blending it with NR. This is definitely associated with the strain-induced crystallizationof NR. The Young’s modulus decreased from N0 to N100, which indicates that the initial stretching of SBRand the blend with the higher SBR content requires higher stress. In the swollen state, there is overallreduction in the magnitude of all the mechanical properties. The tensile behavior of the swollen specimens isgoverned by two types of relaxation mechanisms: the intramolecular motions of segments and the molecularmotions involving the adjustments and shifting of chain entanglements. In the equilibrium swollen state, therubber–solvent interaction is maximum and the rubber–rubber interaction is minimum. This gives rise to theabrupt decrease of the tensile properties of the swollen samples. The mechanical properties of the deswollensamples showed an improvement compared to the unswollen samples. This might be due to increase in theinterchain interaction after a sorption–desorption process.

The mechanical properties of the blends were also intermediate to those of the homopolymers. Theproperties of the swollen samples were largely reduced due to the high rubber–solvent interaction duringswelling.




Table 7.2 Mechanical properties of SBR/NR blends: N0 (SBR100); N30 (SBR70/NR30); N50 (SBR50/NR50);N70 (SBR30/NR70); N100 (NR100) [20].

Secant Modulus (MPa)

Sample System

TensileStrength(MPa)

Elongationat Break

(%)

Young’sModulus

(MPa) M100 M200 M100

Unswollen N0 1.82 391 1.36 0.43 1.05 1.24N30 4.59 862 1.22 0.62 1.01 1.35N50 5.24 869 0.80 0.66 0.98 1.34N70 6.09 1007 0.62 0.64 0.96 1.27N100 6.63 1069 0.35 0.37 0.58 0.79

Swollen N0 0.70 225 0.68 0.43N30 0.90 291 0.49 0.38 0.66N50 1.04 342 0.59 0.36 0.63 0.90N70 0.64 362 0.43 0.27 0.42N100 0.54 390 0.32 0.20 0.32 0.42

Deswollen N0 1.89 398 1.72 0.883N30 4.62 870 1.545 0.749 1.119 1.524N50 5.40 892 1.69 0.814 1.159 1.60N70 6.25 1012 0.603 0.959 1.343N100 6.81 1080 0.976 0.625 1.04 1.40

7.2.1.4 Influence of crosslinking on mechanical properties

The degree of crosslinking has a strong influence on the mechanical properties of rubber/rubber, rub-ber/thermoplastic and other blends where one or both phases could be crosslinked. Semba et al. added dicumylperoxide (DCP) to PLA/PCL binary blend to induce chemical crosslinking, to generate a high performancematerial [21]. The effects of crosslinking on the mechanical properties of polylactic acid/polycaprolactoneblends were studied in detail. There were clear differences at the fracture surfaces obtained after tensile test-ing, as shown in Figure 7.9. Many dropout traces were observed in the sample without DCP, which decreasedwith increasing DCP content. The dropout traces of the samples with 0.2 and 0.3 phr of DCP were very few,which is an indication of the good interfacial adhesion between PLA and PCL phases. These results werein agreement with the observed ultimate strain. Variation of tensile properties with DCP content is shown inFigure 7.10.

7.2.1.5 Influence of Compatibilization on Mechanical Properties

Compatibilization is defined as the process of modification of the interfacial properties of the immiscible blend,leading to the creation of a polymer alloy which is an immiscible polymer blend, having a modified interfaceand/or morphology. When two immiscible polymers are blended without compatibilization, one generallyobtains a mixture with a coarse, unstable morphology coupled with poor interfacial adhesion between thephases. As a result, the blends exhibit inferior physical properties to those of individual polymers. A poorinterfacial adhesion results in an immature stress transfer which cannot prevent cracks initiation at theinterface leading to catastrophic failure. Both theories and experiments support the role of compatibilizer inmultiphase polymer systems. It has been assumed that compatibilizers suppress the droplet coalescence inimmiscible polymer blends by preventing droplets from approaching each other [22–24]. Compatibilization ofblends is broadly classified into two types: Physical compatibilization and reactive compatibilization. Physical




Figure 7.9 Tensile fracture surface of 70/30 blends. (a) DCP0; (b) DCP0.1; (c) DCP0.2; (d) DCP0.3. Reprintedfrom [21]. Copyright (2006) with permission from John Wiley & Sons.

compatibilization includes the addition of pre-synthesized copolymer, block (di-,tri-or tapered) copolymer,graft copolymer, random copolymer, gradient copolymer into the immiscible blend (Figure 7.11).

In reactive compatibilization the copolymer is formed in situ by a chemical reaction during the extrusionprocess during the establishment of the immiscible phase morphology. Reactive compatibilization involves aheterogeneous reaction across a phase boundary. Such a reaction is limited by the interfacial volume availableat this phase boundary.

For incompatible blends containing at least one semi-crystalline component, the final tensile properties aredetermined by two competing factors: the increase in compatibility due to the presence of more crystallinecomponent and the extent of compatibility between the component polymers. The former is the propertydetermining factor at low strain level and the latter determines the properties at high strain level.

The effects of compatibilization on the mechanical behavior of the PP/HDPE blends were studied [26].Stress-strain behavior of uncompatibilized PP/HDPE blends is demonstrated in Figure 7.12. These curveswill also give an approximate indication of the maximum tensile strength (σ m), elongation at break (Eb) andYoung’s modulus (E).

The σ m and Eb showed negative deviation from the additivity line indicating that the blends are highlyincompatible. Interestingly, it is found that Young’s modulus experienced a synergism. Compatibility of the




Figure 7.10 Tensile modulus, strength, and ultimate strain of injection moldings. (a) tensile modulus; (b) tensilestrength; (c) ultimate strain; (•) neat PLA; (©) E70/30; (×) neat PCL. Reprinted from [21]. Copyright (2006)with permission from John Wiley & Sons.

blends does not play a major role in the case of Young’s modulus since it is measured at low strain level.The synergism can be explained in terms of interfacial deformation of the blends. It is believed that duringcrystallization of the matrix, deformation of the dispersed particle occurs and this results in the deformationof the interface in polymer blends. Thus during the crystallization of the blend where PP is the matrix phase,solidification of PP occurs in the presence of HDPE melt which constitutes the dispersed phase. During thisprocess, molten HDPE flows into the region between PP sperulates growing near the interface. This resultsin the deformation of the interface between PP and HDPE, and the deformation ends with completion ofcrystallization sperulates, when all the PP melt is converted to sperulates. The net result is an increase ininterfacial area. However, the formed interfaces are so weak that they can transfer stresses only at very lowstrain levels. This is the reason for the synergism in Young’s modulus, which is measured at low strain levels.




Figure 7.11 Schematic connecting chain at an interface. (a) diblock copolymers; (b) end grafted chains; (c)triblock copolymers; (d) multiply grafted chains; (e) random copolymer. Reprinted from [25]. Copyright (2002)with permission from Springer.

The effect of compatibilization on the σ m of 80/20 blends is given in Figure 7.13. Ethylene propylene dieneterpolymers (EPDM) with three different ethylene /propylene(E/P) ratio were used as the compatibilizers.

The σ m of the blends increased by the incorporation of compatibilizers irrespective of the difference intheir symmetry in terms of monomer fraction. However, it is also seen that the σ m increased with increasein the amount of compatibilizer, reached maximum at 5 wt% of the compatibilizer and beyond this almostleveling off in strength is observed. This is in good agreement with the morphology of the blends, whichrevealed that 5wt% compatibilizer concentration was the interfacial saturation point (CMC) beyond whichno effective compatibilization took place.

Str

ess

(MP

a)

Strain (%)

35

30

25

20

15

10

5

00 2 4 6 8 10 12 14 16 18 20

H0H20H50H80H100

Figure 7.12 Stress-strain behavior of uncompatibilized PP/HDPE blends [26].




Ulti

mat

e te

nsile

str

engt

h (M

Pa)

Weight % of compatibiliser

34

0

0.5E0.6E0.7E32

30

28

265 10 15 20

Figure 7.13 Effect of compatibilization on the tensile strength of HDPE/PP blends (20 parts HDPE).0.5E, 0.6E,0.7E indicate 0.5, 0.6 and 0.7 of ethylene content in Ethylene propylene diene terpolymers respectively [26].

On the other hand, elongation at break (Eb) of the blends increased with compatibilizer concentration.Even though there is no true leveling off in elongation, the rate of increase decreased beyond 5w% of thecompatibilizer concentration. Further, unlike tensile properties, there is difference in performance of differentcompatibilizers.

The Young’s modulus presented in Figure 7.14, on the other hand, shows a different behavior. As thecompatibilizer concentration increased, the Young’s modulus experienced marginal decrease. This is basicallydue to the combined effect of two phenomenon. First, since Young’s modulus is measured at low strain level,as mentioned earlier, compatibility has no significant role. The presence of compatibilizer indeed increased

Figure 7.14 Effect of compatibilization on the Young’s modulus of HDPE/PP blend containing 20 parts HDPE[26].




the compatibility and therefore enhanced most of the mechanical properties of the blends, but did not havean impact on the Young’s modulus. Thus increase in compatibility does not increase the Young’s modulus.

There are two possibilities for the lower values of Young’s modulus in the presence of compatibilizer: (i)the addition of compatibilizer may reduce the crystallinity of the individual polymer and thereby decreasesthe Young’s modulus. This can be ruled out as we have seen compatibilizers have no significant effect in thecrystallization behavior of PP and HDPE; (ii) the compatibilizers may increase the softness of the blends andthis decreases the Young’s modulus. Since the compatibilizer is a flexible polymer, this is a more probablereason for decrease in Young’s modulus [27, 28].

Kim et al. prepared the blends of PBT and EVA by reactive compatibilization of PBT and EVA by MAH(maleic anhydride) [29]. The formation of PBT-g-EVA copolymer as an in situ compatibilizer identified bythe reaction of hydroxyl groups and/or carboxylic groups at the chain ends of PBT and MAH grafted ontoEVA in the presence of DCP (dicumyl peroxide) as an initiator. When EVA is blended with PBT, the in situcompatibilizer, i.e., PBT-g-EVA, has obtained from the reaction of MAH grafted onto EVA and the hydroxylgroups and/or carboxylic groups at the chain ends of PBT. The grafting yield and the gel content duringthe reactive compatibilization processes were higher at higher DCP content. The mechanical properties ofcrosslinked polymeric materials are usually improved with the degree of crosslinking. As a result it wasnoticed that the flexural strength of the PBT-EVA-g-MAH blend is apparently affected by the crosslinkedcomponents of EVA-g-MAH, while the tensile strength is not. The tensile strength decreased but flexuralstrength increased with the increasing gel contents and grafting yield.

7.2.1.6 Influence of Fillers on Mechanical Properties

In recent years, new kinds of polymeric materials are emerging, such as polymer blends reinforced withmicro- and nanofillers such as glass beads, talc, organically modified clays, nanosilica and carbon nanotubes(CNT), which are attracting immense attention because of their remarkable properties.

These new kinds of high-performance materials combine the advantages of polymer blends and the merits ofpolymer nanocomposites. For multi-component systems, one can envision to move beyond simple dispersionand to design strategies that afford the opportunity to selectively reinforce only one of the polymer phases,or even design nanofillers that can promote compatibilized structures of the phases [30–36].

Rahmatpour et al. prepared blend/clay nanocomposites of 50/50 (wt%) NR/SBR [37]. Figure 7.15 showsthe mechanical properties of 50/50 NR/SBR blend/clay nanocomposites. The hardness (shore A) and 100%tensile stress (i.e., stress at 100% elongation) of the nanocomposites are higher than those of clay-free50/50 NR/SBR blend vulcanizate, and also increase with increasing amounts of clay, up to 6 phr, afterwhich they remained almost constant. Increase in the hardness and 100% tensile stress of nanocompositesrelative to the clay-free vulcanizate can be attributed to the layered structure of clay and extremely highinterfacial action between the silicate layers (or stacked layers) and rubber matrix. The tensile strength andtensile strain at break of NR/SBR blend also are improved by introducing the clay into the rubber matrix.In addition, the improvement increased by increasing the amount of clay up to 6 phr and then decreasedslightly in the nanocomposite containing 10 phr clay. It should be noted that all of the mechanical propertiesof nanocomposites improve considerably with an increasing clay content, up to 6 phr, and then remain almostconstant. In other words, maximum improvement in the mechanical properties of nanocomposites can beachieved only when all of the layered silicates are separated into single layers.

Vo and Giannelis have prepared poly (vinylidene fluoride)/nylon-6 (PVDF/N6) blend clay nanocomposite[38]. In their study, they used nanoclay as an alternative means to control interfacial properties. Nanoclaysare an attractive alternative to traditional compatibilizers. This polymer blend/clay system shows improvedstiffness, strength, and toughness. Two different type of clay (Cloisite 30B and Cloisite 20A) and twodifferent mixing stratagies (one batch and sequential compounding) were used for the blend nanocomposite




Tens

ile s

tren

gth

(MP

a)

Tens

ile s

tren

gth

at b

reak

(%

)

Clay amount (phr)

10 700

0 2 4 6 8 10

Tensile strengthTensile strain

7

4

1

8

5

2

9

6

3

012

300

400

500

600

Figure 7.15 Mechanical properties of clay-free 50/50 NR/SBR blend vulcanizate and 50/50 NR/ SBR blend/Na-MMT nanocomposites. Reprinted from [37]. Copyright (2008) with permission from Taylor & Francis.

preparation. They found that toughness is related to the domain size and degree of crystallinity of the domainphase. The domain size is controlled by the PVDF/N6 viscosity ratio and interfacial tension which areultimately determined from the surface modification of the nanoclay and the dispersion of the nanoclayparticles.

To better understand the kinetics of nanoparticle diffusion on blend properties, a series of experimentswere performed where the 30B clay was first compounded with one polymer and then the resulting compositewas compounded with the second polymer. The 30B clay interacts more favorably with N6, so in the blendin which PVDF is blended with 30B first and then with N6 (referred to as (PVDF/30B)/N6), it is expectedthat 30B will migrate to the interface and/or the matrix. For the inverse sequence blend in which 30B iscompounded first with N6 and then with PVDF (referred to as (N6/30B)/PVDF), it is expected that the 30Bnanoparticles will more likely reside in the matrix. TEM images support the overall exfoliation of clay inN6 for both blends (Figure 7.16). Comparison of the TEM images shows a striking difference among theblends’ domain sizes. The PVDF domains in the (PVDF/30B)/N6 blend are ∼110 nm; the domains arenot as small as the one batch blend (∼60 nm) but not as large as the domains of the blend with no clay(∼150 nm). However, the (N6/30B)/PVDF sequential blend exhibits larger domains (∼240 nm) than theblend with no clay. Because of the larger domain size and lack of crystallization suppression in the sequenceblends, the mechanical properties of the sequential blends were poorer than those of the one batch blend(Figure 7.17). The (N6/clay)/PVDF sequence blend with large domains was stiff (E) 2.52 (0.06 GPa) butshowed no improvement in toughness or strength compared to the blend with no clay. Conversely, the(PVDF/clay)/N6 blend was stiffer (E) 2.59 (0.04 GPa), stronger, and tougher compared to the blend with noclay but was inferior to the one batch 30B blend.

7.2.1.7 Modeling of Mechanical Properties

It is well known that mechanical properties might be used to assess the miscibility in polymer blends througha comparison of experimental results and predictions based on various models. Indeed, the mechanicalproperties of polymer blends depend on the intermolecular forces, chain stiffness, and molecular symmetryof the individual polymers used to prepare the blend [39]. Furthermore, according to Willemse et al. [40],




Figure 7.16 TEM micrographs of PVDF/N6 30:70 blend nanocomposites with 5% 30B (a) one batch blend;sequential blend in which (b) clay was compounded first with PVDF [(PVDF/30B)/N6] and (c) clay was com-pounded first with N6 [(N6/30B)/PVDF]. Reprinted from [38]. Copyright (2007) with permission from AmericanChemical Society.

tensile modulus of polymer blends is strongly dependent on the composition and morphology of blends, andtheoretically, it lies between an upper limit given by the parallel model:

M = M1ϕ1 + M2ϕ2 (7.1)

where M is the modulus of the blend, M1 and M2 are the moduli of the components 1 and 2 respectively, ϕ1

and ϕ2 are the volume fraction of the components 1 and 2 respectively. This equation is applicable for modelsin which the components are arranged parallel to the applied stress.




Str

ess

(MP

a)

% Strain

100

0 20 40 60 80 100

20

60

40

80

0120

PVDF:N6 30:70with 1% 30B claywith 5% 30B clay(PVDF+30B) + N6(N6+30B) + PVDF

Figure 7.17 Tensile stress-strain curves for PVDF/N6 blend and 30B modified blends. Reprinted from [38].Copyright (2007) with permission from American Chemical Society.

The lower bound of the modulus holds models in which the components are arranged in series with theapplied stress and the equation is:

1/M = ϕ1/M1 + ϕ2/M2 (7.2)

According to Halpin-Tsai equation:

M1/M = (I + Ai Biϕ2)i(l − Biϕ2) (7.3)

Bi = (M1/M2 − l)/(M1/M2 + Ai ) (7.4)

In the Halpin-Tsai equation, subscripts 1 and 2 refer to the continuous and dispersed phase respectively.The constant Ai is defined by the morphology of the system. For elastomer domains dispersed in a continuoushard matrix, Ai = 0.66.

Moreover, for a cocontinuous system, the modulus has to agree with the Davies model, where the equationis given by:

E1/5 = φ1.E1/5 + φ2.E

1/52 (7.5)

in which Ei and ϕi are the elastic modulus and thevolume fraction of phase i, respectively.The Coran–Patel model [41], which represents a phenomenological adjustment between the parallel and

series models, is

E = φnH (nφs + 1)(EU − EL ) + EL (7.6)

where EU is the upper bound and EL is the lower bound and n is an adjustable parameter related to the changein phase morphology as a function of H.

Kunori and Geil [42] reported that when a strong adhesive force exists between the blend components thedispersed phase will contribute to the strength of the blend. The equation is:

σb = σm(1 − Ad ) + σd Ad (7.7)




where Ad represents the area occupied by the dispersed phase in the transverse cross section. Consideringtwo possible fracture paths in a blend, the equation can be modified as follows, depending on whether thefracture is through the interface or through the matrix. When the fracture is through the interface:

σb = σm(1 − φ2/3d ) + σbφ

2/3d (7.8)

When the fracture is through the matrix:

σb = σm(1 − φd ) + σdφd (7.9)

where σ b, σ m, and σ d are the properties of the blend, matrix phase, and dispersed phase respectively, and ϕd

is the volume fraction of the dispersed phase. Another important model for perfect adhesion is the Kernerequation [43]. According to this:

E = Ee

φd Ed (7 − 5νm)Em

+(8 − 10νm)Ed + φm/15(1 − νm)

φd Em/(7 − 5νm)Em

+(8 − 10νm)Ed + φm/15(1 − νm)

(7.10)

where E, Em, and Ed are the respective properties of the blend, continuous phase, and dispersed phase; ϕd andϕm, the volume fractions of the dispersed and continuous phases; and νm, the Poisson’s ratio of the continuousphase.

Varghese et al. [15] applied the models such as the parallel model, series model and Halpin-Tsai equationto predict the mechanical and viscoelastic behavior of NBR/EVA blends. The applicability of these simplemodels to mechanical and viscoelastic properties is presented in Figure 7.18. In all cases it is seen that theexperimental data are close to the parallel model.

ExperimentalParallelSeriesHelpin–Tsai equation

ExperimentalParallelSeriesHelpin–Tsai equation

Tens

ile s

tren

gth

(MP

a)

15

10

5

0

Volume fraction of NBR

0 10080604020

Sto

rage

mod

ulus

, E′ (

x107 N

/m2 )

Volume fraction of NBR

200

0

100

150

50

010080604020

(a) (b)

Figure 7.18 (a) Applicability of various models on tensile strength of NBR/EVA blends; (b) Applicability ofvarious models on storage modulus of NBR/EVA blends at 10◦C [15].




TE

NS

ILE

ST

RE

NG

H (

MP

a)

VOLUME FRACTION OF NR

7

0

6

5

4

3

2

1

00.3 0.5 0.7 1

Experimental ParallelKunori Kerner

Series

Figure 7.19 Comparison of experimental tensile strength with theoretical values as a function of volume fractionof NR [20].

Mechanical behavior of SBR/NR blends was modeled [20] using various composite models such as theparallel, the series, the Kerner and the Kunori models. Figure 7.19 shows the theoretical and experimentalcurves of the tensile strength values of the SBR/NR blend. In N30 and N50, the experimental values are closeto that of the Kunori model. Therefore, it can be concluded that the fracture propagates through the interfacerather than through the matrix. In N70, the experimental value is close to that of parallel model, Therefore, inN70, it can be concluded that the applied stress distributes equally in two phases.

Veenstra et al. [17] proposed a model to predict the moduli of polymer blends with co-continuous mor-phologies over the complete composition range. This model is obtained by depicting the co-continuousmorphology as three orthogonal bars of the first component embedded in a unit cube where the remainingvolume is occupied by the second component, leading to a series model of parallel parts and a parallelmodel of serial-linked parts. A judicial use of these models results in a perfect description of the moduli ofco-continuous blends as a function of composition.

To visualize co-continuity, the model consists of three orthogonal bars of polymer 1 embedded in a unitcube where the remaining volume is occupied by component 2. Repeating this unit cube in 3D shows thatcomponent 2 has the same framework as component 2, i.e., both the components are interconnected. Relationsfor a series model of parallel parts (Figure 7.20(a) and Eq. (7.11)) and for a parallel model of serial-linkedparts (Figure 7.20(b) and Eq. (7.12)] can be derived [44] as:

Ec = (a4 + 2a3b)E21 + 2(a3b + 3a2b2 + ab3)E1 E2 + (2ab3 + b4)E2

2

(a3 + a2b + 2ab2)E1 + (2a2b + ab2 + b3)E2(7.11)

Ed = a2bE21 + (a3 + 2ab + b3)E1 E2 + ab2 E2

2

bE1 + aE2(7.12)

where a is related to the volume fraction of component 1 by 3a2 2 2a3 = ϕ1 and b is related to the volumefraction of component 2 by b =1− a.




Figure 7.20 Three-dimensional models for the calculation of the moduli of co-continuous polymer blends (aand b). Reprinted from [17]. Copyright (2000) with permission from Elsevier.

They compared the Young’s moduli of co-continuous blends of polyethylene (PE) and polystyrene (PS)to the models as proposed in Eqs. (7.11) and (7.12). Again they used Eq. (7.12) in the region (35–45 vol.%)where the soft phase dominates, Eq. (7.11) in the region (50–80 vol.%) where the stiff phase dominates, andan average between both the equations in the intermediate range where neither phase dominates. The linearinterpolations result in a stepwise plot. The results can be seen in Figure 7.21. The agreement between themoduli that are predicted by the models and the experimental data for the co-continuous PE/PS blends isagain satisfactory.

7.2.2 Dynamic mechanical properties

Dynamic mechanical thermal analysis (DMTA) is another powerful technique to investigate the performanceof polymer blends, as it measures responses of a material to cyclic stress. The investigation of dynamic modulus

5000

0.0

1000

1000.1 0.2 0.3 0.4

Volume fraction PS

Ey

(MP

a)

0.5 0.6 0.7 0.8 0.9 1.0

Figure 7.21 Young’s moduli of co-continuous PE/PS blends (•) and the predictions (full line) using Eqs. (7.11)and (7.12). For comparison, the parallel (—), series (– – –) and Davies (. . . . . .) model are shown. Reprinted from[17]. Copyright (2000) with permission from Elsevier.




and damping behavior over a wide range of temperatures and frequencies has proven to be very useful instudying the structural features of polymer blends and the variation of properties with respect to end useapplications [45–52]. These rely on structure, crystallinity, extent of crosslinking, etc., which in turn dependson the phase morphology of blends. Note that the dynamic mechanical properties are sensitive not only todifferent molecular motions but also to various transitions, relaxation processes, structural heterogeneity andthe morphology of multiphase systems. Further, the dynamic mechanical properties of polymers give mirrorimages of their molecular and morphological features.

Another important aspect is that polymers are viscoelastic materials, which have some of the characteristicsof both viscous liquids and elastic solids. Elastic materials have a capacity to store mechanical energy withno dissipation of energy; on the other hand, a viscous fluid in a non-hydrostatic stress state has a capacityfor dissipating energy but none for storing it. When polymeric materials are deformed, part of the energyis stored as potential energy and part is dissipated as heat. The energy dissipated as heat manifests itself asmechanical damping or internal friction. Therefore the interpretations of these properties at molecular levelare of great scientific and practical importance in understanding the mechanical behavior of polymers.

Dynamic mechanical thermal analysis helps to measure the glass transition temperature of polymers. Inaddition one can obtain an idea about the storage (dynamic) modulus, loss modulus, and damping behavior(internal friction). The Tg of polymers is a very important parameter since there are profound changes in thephysical properties of the polymers –heat capacity, thermal expansion coefficient and modulus occur at thistemperature. It should be noted that Tg of the polymer is accompanied by a sharp decrease in stiffness. Thestress relaxation modulus decreases and the creep compliance increases by about three orders of magnitudein glass transition region. The loss moduli and loss compliances exhibit a maximum in the glass transition asdoes tan δ. Any type of interaction between the polymers will give rise to a shift in this maximum and thusdynamic mechanical properties will provide an idea about the extent of interactions between the componentpolymers in a polymer blend.

The effect of different compatibilizers on the dynamic mechanical properties of PS/polybutadiene blendswas studied by Joseph [53]. The author reported that storage modulus, loss modulus and tan δ underwentdramatic change in the presence of compatibilizers. For LLDPE/EVA blends, it has been found that compati-bilization increased the storage modulus of the system which is due to the dispersion of EVA domains in theLLDPE matrix providing an increased interfacial interaction [54].

Figure 7.22 demonstrates the effect of blend ratio on the storage modulus (E′) of PP/HDPE blends [26].PP exhibits the maximum and HDPE the minimum storage moduli in the whole temperature range. The E′

of all blend exhibits intermediate values. As the weight % of HDPE in the blend increases, storage modulusincreases. However, it is interesting to compare the Young’s modulus with storage modulus since both aremeasured under tension. However, the E′ values offer apparent conflict with the Young’s moduli and areconsiderably higher than the Young’s moduli. At the same time there is no synergism in the E′ values.

The effect of blend ratio on the Tg of PP are given in Figure 7.23. Separate peaks suggest that the HDPE/PPblend is highly incompatible at all combinations.

The vulcanization of the rubbery phase during mixing has been investigated as a way to improve the physicalproperties of several thermoplastic elastomers based on rubber/plastic blends. The change in morphologythat occurs during dynamic vulcanization is schematically represented in Figure 7.24. During dynamicvulcanization, a co-continuous morphology may be transferred to a matrix and dispersed-phase morphology,there may be some possibility of phase inversion, or the crosslinked rubber phase may become finely anduniformly dispersed in the plastic matrix. During the process of dynamic vulcanization, the viscosity ofthe rubber phase increases because of crosslinking, and the rubber domains can no longer be sufficientlydeformed by the local shear stress and are eventually broken down into small droplets.

The variations of G′, G′′ and tan δ of 70/30 HDPE/ EVA blends crosslinked with 0.5 and 1.5% DCP (dicumylperoxide) were studied by John et al. [55]. The addition of peroxide to 70/30 HDPE/EVA blends leads to




Figure 7.22 The variation of storage modulus of PP, HDPE and their blends as a function of temperature [26].

an increase in G′ and G′′(Figure 7.25). This increase is more pronounced at the higher DCP contents. Attemperatures greater than 100◦C, 1.5% DCP leads to a decrease in G′′. The α- and γ -relaxation temperaturesof HDPE are not influenced significantly by the addition of DCP. The relaxation of EVA at about –20◦C isslightly changed by the addition of 0.5% DCP. The addition of 1.5% DCP leads to significant increase in therelaxation temperature of EVA by 3.6 K, which indicates the predominant crosslinking of the EVA phase.In addition, the intensity of the tan δ peak corresponding to the EVA content increases with the increasein the DCP content. In this case also, the blends show the presence of two peaks corresponding to Tg’s ofHDPE and EVA, which indicates the immiscibility of the system. The peak widths at the half-heights of the

Figure 7.23 Effect of blend ratio on the variation of tan δ as a function of temperature in HDPE/PP blend [26].




Figure 7.24 Schematic representation of the dynamic vulcanization [55].

Sto

rage

mod

ulus

G’(P

a)

Temperature (°C)

1010

–150

HDPE/EVA70/3070/30 + 0.5% DCP70/30 + 1.5% DCP

109

108

107

–100 –50 0 50 100 150

Figure 7.25 Dependence of G′ on the temperature of 70/30 HDPE/EVA dynamically crosslinked blends [55].




Table 7.3 Storage modulus of LDPE, POE, LDPE/POE blends, and LDPE/POE/claynanocomposites [56].

Storage modulus/GPa

Sample No. −140◦C −100◦C −50◦C 0◦C 25◦C

1 3.57 2.40 1.67 0.65 0.3415 3.35 2.38 1.72 0.65 0.335 3.51 2.35 1.55 0.36 0.198 3.73 2.49 1.68 0.39 0.2014 3.44 2.40 1.61 0.36 0.1913 4.29 2.90 2.02 0.54 0.2811 3.37 2.23 1.32 0.17 0.0912 3.74 2.50 1.48 0.20 0.109 2.51 1.66 0.85 0.06 0.0310 2.99 2.02 1.08 0.08 0.042 3.49 2.12 0.84 0.03 0.0216 4.08 2.78 1.18 0.04 0.02

70/30 HDPE/EVA blend and the 70/30 HDPE/EVA blends crosslinked with 0.5 and 1.5% DCP were 31, 33,and 37◦C, respectively. This increase in the peak width or broadening of peaks upon crosslinking indicatesthat dynamic crosslinking promotes interfacial bonding between the phases, which may arise because of thecrosslinking of HDPE and EVA phases.

Dynamic mechanical properties of low-density polyethylene (LDPE)/ethylene–octene copolymer(POE)/organomontmorillonite (OMMT) nanocomposites were studied by Baghaei et al. [56]. The storagemodulus of the LDPE/POE blends and nanocomposites at different temperature regions are shown in Table7.3. As can be seen, the storage modulus of the nanocomposites are higher than their neat blends, indicatingthe reinforcing role of clay. OMMT anchors at different positions in the matrix, thus restricting the movementof the chains.

As seen from Table 7.3, at temperatures higher than glass transition temperature (Tg) of POE (–50 ◦C)storage modulus of the POE/OMMT (95/5) nanocomposite is about 50% higher than that of the neat POE.The enhancement of storage modulus strongly depends on the aspect ratio of the dispersed clay layers and theintercalation of the polymer chains inside the clay matrix. When the polymer matrix is reinforced with a rigidfiller the polymer interface adjacent to the clay particle is highly restrained mechanically. Active surface areaof the filler increases due to the intercalation of the polymer chains inside the clay galleries. Polymer chainsinside the clay galleries are immobilized and the effective immobilization of these chains is responsible forthe enhancement of the hydrodynamic storage modulus

7.2.2.1 Modeling of Dynamic Mechanical Properties

Kim et al. applied parallel-series model, Davis model and Coran–Patel model to methyl acrylate (MA) and2-acrylamido-2-methyl propane sulfonic acid (AMPS abbreviated as AP) modified polyacrylonitrile/celluloseacetate (MA–PAN–CA and AP-PAN-CA) blends [57]. MA–PAN–CA blends (Figure 7.26) shows negativedeviation from the Davies model, whereas AP–PAN–CA blends (Figure 7.27) are relatively well fitted withthe model. These results imply that interfacial interactions are poor in MA–PAN–CA blends and relativelygood in AP–PAN–CA blends, and the results are consistent with dynamic mechanical data and morphology.




0

10.2

9.6

9.9

9.3

9.0

n = 1n = 2n = 3n = 4n = 5

20 40

CA content (wt%)

Log

E(d

yne/

cm2 )

60 80 100

Figure 7.26 Complex moduli of MA–PAN–CA blends at 80◦C, and comparison with upper–lower bound model(solid line), Davis model (bold solid line), and Coran–Patel model (dashed line). Reprinted from [56]. Copyright(2009) with permission from Springer.

For MA–PAN– CA blends, the 80/20 blend is fitted with n = 3 of the Coran–Patel model (Eq 7.6), 40/60 and60/40 blends with n = 3 – 4, and the 20/80 blend with n = 5, and the experimental values are generally wellfitted with n = 3 – 4. However, in the AP–PAN–CA blend, experimental values are well fitted with n = 1 – 2.(n − 1)/n indicates the center of the H range where phase inversion or transition occurs. Thus, it can be notedthat phase inversion takes place at H = 0.67 – 0.75 (MA–PAN–CA) and 0.5 – 0.67 (AP–PAN–CA blend).These results agree well with morphology studies.

7.2.3 Impact Properties

Extensive research and development work have been carried out to formulate polymers with high impactresistance. Both rubbers and engineering thermoplastics are incorporated into brittle plastics to improve their

n = 1n = 2n = 3n = 4

10.2

Log

E(d

yne/

cm2 )

CA content (wt%)

10.0

9.8

9.6

9.40 10080604020

Figure 7.27 Complex moduli of AP–PAN–CA blends at 80◦C, and comparison with upper–lower bound model(solid line), Davis model (bold solid line), and Coran–Patel model (dashed line). Reprinted from [56]. Copyright(2009) with permission from Springer.




impact properties [58–60]. Impact modification by rubber toughening involves the incorporation of smallamounts of rubber (mostly between 3 to 20 vol%) in rigid polymer materials such as glassy thermoplastics,semicrystalline thermoplastic, and thermosets in order to enhance their fracture resistance. The impact resis-tance of rubber-toughened polymers is strongly dependent upon the concentration, size and size distributionof the rubber particles [61].

The role of particle size distribution has been studied by Wu [62]. According to Wu, a wide particlesize distribution is disadvantageous for the impact behavior of rubber-modified blends with pseudoductilematrices because the interparticle distance increases with the increasing size of polydispersity at a give rubberfraction. According to Bucknall and Paul, in high-performance blends, toughness reaches a maximum atparticle diameters of 0.2–0.4 μm [63].

The critical interparticle distance (IPD) is related to the size of the rubber particles and the rubber volumefraction, ϕr, by the following relation.

IPD = d[k(/6ϕr )1/3l] (7.13)

where the parameter k is a measure of the packing arrangement of a particular lattice and d is the averagediameter of the dispersed rubber particle.

A bimodal distribution can lead to an enhancement of the toughness. With respect to the role of interfacialadhesion, a distinction has been made between matrices deforming by multiple shear yielding and thosedeforming by multiple crazing. For matrices deformed by multiple crazes, good adhesion between the rubberparticle and matrix is required because the rubber particles must act as effective craze stoppers. For matricesdeformed by shear yielding as a result of internal rubber cavitation, a good interfacial strength is not required.The particle size required for the cavitation, d, is given by the following equation:

D = 12(γr + sc)/Kr4/3 (7.14)

where is the volume strain, Kr is the rubber bulk modulus, sc is the surface energy per unit area, and γ r

is contribution from van der Waals surface tension.Impact properties measurements of PA 6/EPR showed that notched impact strength curves are strongly

influenced by the particle size, as shown in Figure 7.28 [8].Borggreve and Gaymans studied the effect of the coupling agent on the impact behavior of nylon 6-

EPDM (ethylene propylene diene monomer) rubber. The rubber was grafted with various amounts of maleicanhydride (MA) with the aid of a peroxide. The MA grafted onto the rubber was found to react with thenylon during the blending process. With the MA-grafted rubbers, a much finer dispersion could be obtained.However, the concentration of the coupling agent, within the range 0.13 to 0.89 wt% grafted onto the rubber,has hardly any influence on either the dispersion process or the impact behavior of the blends [64].

The influence of mixing conditions on mechanical properties of low density polyethylene–polystyreneblends were studied by Vasilenko et al. [65]. LDPE–PS blends were obtained by two methods: (a) bymechanical mixing in a melt at 160◦C in a closed mixer, and (b) by the (high temperature shear deformation)HTSD method in a rotary disperser. For PS, the impact strength is one of the most important mechanicalcharacteristics. It can be improved by introducing LDPE. Figure 7.29 shows that the impact strength of theLDPE–PS blends prepared by the HTSD method exceeds by a factor of 2.0–2.5 the impact strength of thesamples obtained by mixing in the melt.

Semba et al. added dicumyl peroxide (DCP) to PLA/PCL binary blend to induce chemical crosslinking,to generate a high performance material [21]. The effect of crosslinking on impact properties of polylacticacid/polycaprolactone blends were studied in detail. The impact strength of PLA maintained a constant valueat all DCP contents in both tests. On the other hand, the impact strength of PLA/PCL (70/30) showed a higher




Figure 7.28 Noted Izod measurements as a function of the weight average particle size for PA6/EPR (20 vol%)[8].

value, which increased with increasing DCP content at low DCP concentrations. In the case of the normalIzod impact test, the PLA/PCL containing 0.3 phr of DCP showed a value that was 2.5 times the value of thePLA. In the case of the reverse Izod impact test, the impact strength of both samples showed similar valuesfrom 0 to 0.2 phr of DCP contents. On the other hand, PLA/PCL blend containing 0.3 phr showed a highervalue that was 2.5 times those of the other materials. Based on those observations, it is obvious that DCPaddition significantly alters the impact properties of this blend system.

50

Rel

ativ

e im

pact

str

engt

h, k

J/m

2

LDPE content, wt %

0

25

025 50 75

2

1

Figure 7.29 Relative impact strength of the LDPE–PS blend vs. the content of LDPE with respect to the impactstrength of PS: (1) HTSD mixing and (2) mixing in the melt. Reprinted from [65]. Copyright (2009) with permissionfrom Springer.




1200

IZO

D IM

PAC

T S

TR

EN

GT

H (

J/m

)

% WEIGHT

0NYLON 6 PP

1000

800

600

400

200

020 40 60

20% SEBS-g-MA

20% EPR-g-MA

0% Rubber

80 100

Figure 7.30 Izod impact strength of nylon 6/polypropylcnc blends modified with 20% EPR-g-MA or SEBS-g-MA;the composition shown on the abscissa is the percentage of polypropylene relative to nylon 6 in the blend ona rubber-free basis. The dashed lines correspond to the properties of unmodified binary blends. Reprinted from[66]. Copyright (1995) with permission from Elsevier.

Paul and co-workers studied the effect of the two maleated rubbers EPR (ethylene-propylene ran-dom copolymer)-g-MA and SEBS (styrene-ethylene/butylenestyrene triblock copolymer)-g-MA, containingequivalent amounts of maleic anhydride, on the mechanical properties of nylon 6/PP blends [66]. Figure 7.30shows that the Izod impact strength is greatly improved by addition of 20% by weight of the maleated rubbers,equivalent to 25.1% and 23.8% by volume of EPR-g-MA and SEBS-g-MA, respectively. The improvement inimpact strength is relatively independent of the type of rubber used. High impact strengths were obtained inthe range 0–50% of polypropylene. Beyond 50% polypropylene, the impact strength is lower: this reductionappears to occur at the point where the polypropylene becomes the continuous phase. The impact strengthshows a maximum at about 20–25% polypropylene for both rubbers.

Toughening of thermoset for improvement of crack resistance has been the subject of intense researchinterest during the last two decades. Reactive liquid rubbers and functionally terminated engineering ther-moplastics are the widely used toughening agents for the epoxy thermoset. Carboxyl-terminated butadieneacrylonitrile (CTBN) copolymer has been used to modify the aromatic amine cured epoxy resin and anhydridecured epoxy [67, 68]. Appreciable improvements in impact strength were observed in the prepared blendsystems (Figure 7.31).

Impact behavior of the cured epoxy could be explained based on the two-phase nature of the system.According to Bucknall [61] the rubber particles were considered to bridge the crack as it propagates throughthe material. Thus, the rubber particles were able to prevent the crack growing to a catastrophic size. Theincrease in toughness was due to the amount of elastic energy stored in the rubber particles during stretching.Thus, the deformation of the rubber particles in the matrix seemed to be responsible for the enhanced stresstransfer and hence impact resistance. Shear yielding of the matrix was another reasonable mechanism thatmight be operating. According to Newman and Strella [69] the principle function of the rubber particle wasto produce sufficient triaxial tension in the matrix so as to increase the local free volume and hence enableextensive shear yielding of the matrix. Thus, crack building of rubber particles along with shear yielding wasthe main toughening mechanism and enhancement of impact behavior.

7.2.4 Nanostructured Polymer Blends

The idealized morphology of nanostructured polymer blend systems is characterized by the molecular leveldispersion of the phases that leads to a considerable enhancement in the mechanical properties, especially




Impa

ct s

tren

gth,

J x

cm

2

Toug

hnes

s, J

CTBN, %

50300

250

200

150

100

20

30

40

0 252015105

Figure 7.31 Variation of impact strength and toughness with CTBN content in epoxy/CTBN blend cured witharomatic amine. Reprinted from [67]. Copyright (2007) with permission from Elsevier.

the modulus. In recent years, several strategies have been developed from well-defined and predictable mul-ticomponent polymer structures with phase separation at nanoscale [70]. The most straightforward approachis to use linear block copolymers with two components, A and B. Significantly different morphologies andmechanical properties were generated when replacing a mixture of two almost immiscible linear polymers(PS and PMMA) by the corresponding block copolymer (PS-b-PMMA), to modify an epoxy resin. As shownin Figure 7.32, the morphology obtained for the BC/E (block copolymer-epoxy) blend was quite differentthan the PS/PMMA/E blends. The morphology of BC/E blend is characterized by the presence of bicon-tinuous phases, a fact which is highly desirable when the aim is to increase the fracture toughness of thethermoplastic/thermoset blend.

Phase separation of the more miscible block induced by polymerization led to the generation of a bicon-tinuous thermoplastic/thermoset structure exhibiting the desired decrease in yield stress which is necessary

Figure 7.32 SEM micrographs of (a) PS/PMMA/E blends; (b) BC/E blends. Reprinted from [71]. Copyright(2003) with permission from American Chemical Society.




Table 7.4 Elastic modulus (E) and yield stress (σ Y) of the neat epoxy matrixand different blends [60].

Blend E(GPa) δγ (MPa)

Neat E 2.65±0.06 103.3±0.4PS/E 2.54±0.01 97.1±1.2PMMA/E 2.70±0.05 98.0±0.5PS/PMMA/E 2.64±0.01 97.2±1.0BC/E 2.14±0.02 90.0±1.8

for toughening purposes. Values of the elastic modulus (E) and yield stress (σ Y) of the neat epoxy matrix anddifferent blends are shown in Table 7.4.

Ritzenthaler et al. studied the mechanical properties of nanostructured blends of polystyrene-block-polybutadiene-block-poly-(methyl methacrylate) (SBM) copolymer triblock and epoxy as a function ofcomposition and concentration of SBM [71]. The addition of SBM triblocks can then be expected to be apowerful way of improving the poor fracture resistance of epoxy thermosets. Two SBM copolymers wereused: S27

22B9M69-SB21 and S1412B18M70-SB10 (S27

22B9M69-SB21 represent a copolymer as received andcomposed of S27

22B9M69 pure triblock copolymer and 21 wt % of SB diblock. The numbers 22, 9, and69 represent the weight of the respective PS, PB, and PMMA blocks; 27 is the molar mass of the PS block inkg mol-1). Figure 7.34 shows the KIc values obtained for different concentrations of SBM up to 50 wt %. Itis worth noting that the continuous matrix always remains the epoxy-diamine component whatever the SBMconcentration in the investigated range. This constitutes a major difference and advantage compared to theclassical use of low molar mass reactive rubbers, for which phase inversion occurs when the concentration ishigher than 20 wt %. For both SBM copolymers, the toughness very significantly increases with the triblocks

(b)(a)

Figure 7.33 Transmission electron micrographs DGEBA-MCDEA epoxy systems containing 30 wt %. (a)S27

22B9M69-SB21; (b)S1412B18M70-SB10. Reprinted from [71]. Copyright (2003) with permission from American

Chemical Society.




2.5

2

1.5

1

0.5

00 20

KIC

(M

Pa.

m1/

2 )

10 30

SBM wt %

40 50 60 70

Figure 7.34 Influence of as-received SBM concentrations and compositions on material toughness: (*) and neatDGEBA-MCDEA, (•) S27

22B9M69-SB21, (�) S1412B18M70-SB10 as toughner agents at different concentrations.

Reprinted from [71]. Copyright (2003) with permission from American Chemical Society.

concentration. Until 10 wt % of SBM, the toughness is slightly higher for the blends based on the SBMcontaining the longer PB block. Alternatively, for higher concentrations, blends based on S27

22B9M69-SB21exhibit higher toughness, suggesting that the ‘raspberry’ morphology (Figure 7.33a) is more efficient inimproving mechanical properties than the ‘onion’(Figure 7.33b) morphology.

Recently, Thio et al. have reported the behavior of poly-(ethylene oxide)-b-poly (hexylene oxide)(PEO–PHO) diblocks in phenol novolac cured BADGE [72]. The highly non-polar PHO block is immisciblewith BADGE even at low molecular weights. Wormlike micelles were generated by mixing a vesicle formingdiblock (9 wt% PEO) and a spherical micelle forming diblock (44 wt% PEO).Here the wormlike morphology,shown in Figure 7.35, was found to give the best improvement in K1c (∼6×), but vesicles were found to givegreater improvements than spherical micelles (∼3.5× and ∼1.75×, respectively).

7.3 Interpenetrating Polymer Networks (IPNs)

According to Sperling, the IPNs are defined as a combination of two or more polymers in network form thatare synthesized in juxtaposition [73]. For thermoset IPNs, the related networks are chemically crosslinked.The thickness of the network strands varies from 10 to 1000 nm [74]. The IPN structure offers uniquepossibilities for composites with respect to fiber to matrix adhesion as schematically depicted in Figure 7.36.

Thermoset IPNs are prepared sequentially or simultaneously. In the sequential case, the first crosslinkedpolymer network is swollen by the monomer of the second polymer that is polymerized and/or crosslinkedafterward. The simultaneous IPNS are prepared by crosslinking of two or more monomer systems at thesame time. The related reactions should be combination of free-radical induced polymerization and ionichomo-polymerization reactions. The cross-reactions between the crosslinked networks result in grafted IPNs.

Thermoset IPNS have been mostly used as tough, impact modified polymers, matrices in compositesand as sound and vibration damping materials [73, 76, 77]. Several reports are now available on thermoset




Figure 7.35 TEM image showing wormlike micelles of PEO–PHO in a BADGE+ PN system. Stained with RuO4.Reprinted from [72]. Copyright (2006) with permission from American Chemical Society.

IPNs, usually chain and step polymerization crosslinking methods are combined. As mentioned before, chainpolymerization covers free-radical induced crosslinking of polymers and monomers bearing two or moredouble bonds along their chains. This is the case with unsaturated polyester (UPs) with VE resins, andcrosslinking of vinyl compounds possessing one double bond with two or multifunctional monomers ofsimilar structure. Table 7.5 lists the mechanical and fracture mechanical properties of VE/EP hybrids, the EPphase of which was amine cured.

As expected, all IPNS are very sensitive to post curing, which results in a substantial increase in the Tg. Thisis associated with decrease in fracture energy (Gc). It was found that VE/VP hybrids, in which the idealizednetwork of the EP phase is composed of aliphatic and cycloaliphatic units, show outstanding toughness values[78].

Figure 7.36 Scheme of the intermittent fiber bonding achieved via matrix organization (i.e., IPN structure).Reprinted from [75]. Copyright (2003) with permission from Springer.




Table 7.5 Tensile mechanical properties, fracture energy (Gc), and glass transition temperature (Tg) of VE/EP(1:1) systems the EP component of which was cured by polyaddition reaction with diamines [78].

Composition

Maximumcuring

temperature(◦C)

Tensilestrength(MPa)

Ultimatetensile

Strain (%)

Young’smodulus

(MPa)Gc

(kJ/m2)Tg

(◦C)

VE/AI-EP+Cal-Am 150 50 5.5 2600 3.7 81200 60 4.5 2700 1.1 128

VE/Cal-EP+Al-Am 150 14 38.5 700 7.3 57200 46 5.1 2900 5.2 87

VE 150 48 2.8 3200 0.54 157200 52 2.1 3400 0.45 160

Kim et al. evaluated tensile and shear moduli of polyurethane-poly (methy1 methacrylate) interpenetratingpolymer networks (IPNs) [79]. The strain rate dependence of tensile modulus is shown in Figure 7.37.

The modulus-strain rate curves appear to lie in three groups. The low modulus group includes elastomericsamples where the polyurethane phase is continuous, the medium modulus group includes leathery sampleswhere the phases are in the inversion process (each phase is locally connected), and the high modulus groupincludes glassy samples where the poly(methy1 methacrylate) phase is continuous. The 75/25% polyurethane-poly (methy1 methacrylate) IPN shows little modulus reinforcement over the pure polyurethane samples.

Shear storage modulus, G′, vs. temperature plot (Figure 7.38) for polyurethane-poly (methy1methacrylate)IPNs also reveals three general group. The 75/25% and 85/15% polyurethane-poly(methy1 methacrylate)IPN’s show a slightly reduced modulus compared with the pure polyurethane sample at room temperature.

1010

UC75MC25UC60MC40UC50MC50UC40MC60UC25MC75UC15MC85MC100

UC100

109

108

0.001

Youn

g’s

(dyn

e/cm

2 )

Strain Rate (min−1)0.01 0.1 1.0

Figure 7.37 Young’s modulus vs. strain rate for the polyurethanepoly(methy1 methacrylate) IPN’s at 23◦C.Reprinted from [79]. Copyright (1977) with permission from American Chemical Society.




1010

109

108

107

106

She

ar S

tora

ge M

odul

us G

′ (dy

ne/c

m2 )

Temperature (°C)

1107030−10−50−90

UC40MC60UC25MC75UC15MC85MC100

UC75MC25UC60MC40UC50MC50

UC100UC85MC15

Figure 7.38 Shear storage modulus, G’, vs. temperature for the polyurethane-poly(methy1 methacrylate) IPNs.Reprinted from [79]. Copyright (1977) with permission from American Chemical Society.

Modulus studies revealed that IPN has the lowest modulus at room temperature. This is probably dueto crystallization of the linear polyurethane. The above comparison seems to indicate that the reduction ofmodulus is related to the interaction and/or interpenetration of the phases. One possible explanation wouldbe that the partial crystallinity in the polyurethane phase (shown earlier to be present, even in the crosslinkedpolymer) was reduced due to the interpenetration of poly (methy1methacrylate) chains into the polyurethanephase. In the pseudo-IPNs and linear polyblend, the crystallinity of the polyurethane phase was not reducedsince there was no interpenetration (interlocking of chains), thus yielding an increased modulus over thepure polyurethane. The increase in modulus in the temperature range of –10 to 30◦C is due, most likely,to crystallization of the linear polyurethane caused by the slow heating. Moreover, the shift in Tg to highertemperature (most prominent in the full IPNs) also indicates interpenetration.

In another study by Huelck et al., the physical and mechanical properties of poly (ethy1 acrylate)-poly(styrene-co-methyl methacrylate), PEAP(S-co-MMA), interpenetrating polymer networks (IPN’s) have beeninvestigated [80]. Stress–strain and tensile data show that the work to break as well as the actual tensilevalues of the samples steadily increase as the amount of plastic component is increased in the elastomer-richmaterials.

Mechanical properties of the IPNs based on PI/PMMA were correlated to the morphology of the system[81]. Tensile strength of the IPN samples are plotted as a function of the overall PMMA content as shownin Figure 7.39. The IPNs showed a gradual change from a rubbery to a plastic nature with the increasein concentration of PMMA. Tensile strength and modulus were strongly dependent upon the morphologyand showed a dramatic change near the 50/50 composition where the system developed a nanostructuredmorphology (Figure 7.40).

In particular, the tensile strength and modulus showed a lower value in the sea-island region, and a distinctincrease when the morphology changed to compact, ordered and smaller PMMA domains having size lessthan 90 nm and an immediate decrease on the appearance of dual phase morphology. When the PMMAconcentration was higher than 30 wt%, tensile strength increased significantly. Elongation at break was foundto decrease with the increasing PMMA content due to the low elongation and brittleness of the PMMA phase.These results suggest that the nanostructured morphology obtained around 50/50 composition was the keyfactor that contributed to the mechanical performance of these IPNs.




25

20

15

10

5

00 10 20 30

Weight percentage of PMMA

Tens

ile s

tren

gth

MP

a

40 50 60

Figure 7.39 Effect of wt % of PMMA on tensile strength of full IPNs [81].

The introduction of crosslinks in PMMA phase was found to improve the mechanical performance of thematerial. This was due to the higher degree of interpenetration between the PI and PMMA chains during IPNformation. Thus, some degree of enforced miscibility was obtained, which enabled the material to have goodmechanical properties. Higher amount of PMMA (more than 55%) was found to reduce the tensile strength.Effect of crosslink density of PI on the mechanical properties was also studied by preparing IPNs withhighly crosslinked PI samples. A considerable decrease in tensile strength was observed with the increase incrosslinking degree of the PI phase. However, the effect of crosslink density of PMMA on the mechanicalperformance of full IPNs was different from that of the PI phase. The lower crosslink in PMMA favors the

Figure 7.40 TEM image of IPN s with nanometre-sized morphology [81].




Field (G)

3260 3280 3300 3320 3340

193K

ss

ff

293K

303K

333K

353K

383K

413K

Figure 7.41 ESR spectra of probe measured as a function of temperature in semi IPN based on PI and PMMAwith 50/50 composition [81].

mechanical properties up to a certain extent; however, higher crosslinking is not desirable. The combinationof 2% crosslinker concentration in PI and 4% in PMMA in PI/PMMA IPN led to the development ofnanostructured morphology and the said IPN performed better in mechanical measurements.

For a given composition, the elongation at break decreased with the increase of crosslinking in the PI phase.This was due to the reduced chain flexibility of the PI matrix with higher crosslinking. Also crosslinkingincreased the PMMA entanglement density and interpenetration between the phases. The PMMA networkdensity was found to have only little influence on the elongation at break of full IPNs. Young’s modulus wasalso found to increase with the increase in crosslinking of the rubber phase. Also crosslink density of thePMMA phase was found to have a pronounced effect in the Young’s modulus of full IPNs. The 6% crosslinkerconcentration in PMMA showed the highest value among the studied samples. It was found that shore Ahardness increased with the rise in PMMA content.

Moreover, main chain and segmental dynamics of PI and PMMA chains in semi IPNs with differentcomposition, crosslink density, and molecular weight were studied over a wide range of temperatures, usingthe ESR spin probe technique and dynamic mechanical analysis and confirmed the dramatic changes in themotional behavior of both polymers due to the molecular level interpenetration between two polymer chainsin synthesised semi IPN [82]. At lower temperatures, the ESR spectra approached the rigid limit spectrum,and as the temperature increased, the spectral lines got narrower and outer peaks shifted inward. In the rangeof 293 K-373 K, complex spectra were obtained, an indication of the presence of motionally distinct regionsin semi IPNs (Figure 7.41).

7.3.1 Modeling of Mechanical Properties of IPNS

Further development and optimization of different mechanical properties of IPNs requires that experimentalresults be complemented by theoretical studies and computer simulations as they have the advantage ofproviding insights into the role of individual factors, and to rationally elucidate their significance.Various




38

ParallelSeriesHelpsin-TaiTakayanakiKunori 1Kunori 2

Experimentalx

x

x

x

x

3634323028262422

Tens

ile s

tren

gth,

MP

a

wt-% of PMMA

201816141210

8642

0 10 20 30 40 50 60

Figure 7.42 Model fitting–tensile strength of IPNs based on PI/PMMA [81].

models, such as the parallel model, the series model, the Halpin–Tsai equation and the Takayanaki modelwere used to study the mechanical behavior of the IPNs. Kunori and Geil models were used to determinefracture paths (matrix or interface).

John et al. applied the Kunori and Geil models to IPNs based on PI/PMMA. [81]. Figures 7.42 and 7.43show the theoretical and experimental curves of the tensile strength and Young’s modulus values, respectively,

Parallel

Series

Halpsin-Tai

Takayanaki

Kunori 1

Kunori 2

Experimental

900

800

700

600

500

400

300

200

100

0

0 10 20Wt-% of PMMA

Mod

ulus

, MP

a

30 40 50 60

Figure 7.43 Model fitting–Young’s modulus IPNs based on PI/PMMA [81].




BLEND

GRAFT (27)

NORM. IPN

INV. IPN

109

1010

RNP

PINR

PEAB

E′ (

DY

NE

/CM

2 ) A

T 2

5 °C

WT% PS

108

107

50 1000

Figure 7.44 Storage moduli at 25◦C for PEAB-PS IPNs. R stands for rubber and P stands for plastic. The dottedline indicates the theoretical value of the storage modulus for two continuous phases, based on the Bauer model.Reprinted from [80]. Copyright (1972) with permission from American Chemical Society.

for the six models. For both tensile strength and modulus, parallel model fits more closely to the experimentalvalues up to 50 wt% of PMMA and then the values deviate. This may be because both the PI and PMMAphases have a co-continuous morphology above 55 wt% PMMA, rather than having a matrix-dispersed phasemorphology. The theoretical curve of Kunori (equation for fracture through matrix) comes closest to theexperimental curve compared to the other models. Therefore, it may be concluded that the fracture path isthrough matrix rather than through the interface. The Kunori values superimpose over parallel model, so linescannot be distinguished in the figures.

The modulus-composition behavior of the PEA/S-co-MMA IPNs was analyzed with some of the theoreticalequations based on mechanical models [80]. Employing the theory of Bauer et al., which is essentially amodification of the earlier mechanical models of Takayanagi, they judged the relative continuity of the twophases [83]. Figures 7.44 and 7.45 present storage modulus data as functions of composition for PEAB-PSand PEAB-PMMA IPNs, respectively. Also included are the solution blended material, a graft copolymerand data from a mixed latex. The upper solid lines represent a continuous plastic phase, while the lower solidlines represent a continuous elastomer phase. The dotted lines indicate the result obtained for two equallycontinuous phases, all according to theory. Allowing for the relative oversimplification of the mechanicalmodel, the results tend to indicate that both phases do in fact exhibit some degree of continuity, especially thePEA-PMMA compositions. A possible exception is the 75/25 PEAB-PS point in Figure 7.44, which suggestsa continuous elastomeric phase. It should be noted, however, that the theoretical curves were derived foruncrosslinked polyblends.




INV. IPNNORM. IPN

109

1010

RNP

PINR

PEAB

E’ (

DY

NE

/CM

2 ) A

T 2

5 °C

WT% PMMA

108

107

50 1000

GRAFT (26)

BLEND (28)MIX LATX (25)CHCI3 (26)

Figure 7.45 Storage modulus at 25” for PEAB-PMMA IPN’s.Within the limits of the Bauer theory, all such blendsand IPNs appear to have two continuous phases. Reprinted from [80]. Copyright (1972) with permission fromAmerican Chemical Society.

7.4 Polymer Gels

A polymer gel is a crosslinked polymer network swollen in a liquid medium. Its properties depend stronglyon the interaction of these two components. Polymer hydrogels can be divided into two main classes, i.e.,chemically crosslinked hydrogels, which are composed of polymer networks with covalent bonding, andphysically crosslinked hydrogels, which are composed of physical networks with noncovalent interactions.To date, hydrogels of poly(N-alkylacrylamides) such as poly(N-isopropylacrylamide) (PNIPA) and poly(N.N-dimethylacrylamide) (PDMAA), which have attracted extensive attention as water-absorbing, soft and stimuli-sensitive materials, were all prepared by chemical crosslinking reactions using an organic crosslinker such asmethylene bis (acrylamide) (BIS) [84–99].

However. because of the random nature of the crosslinking reactions produced by a large number oforganic crosslinkers, the conventional chemically crosslinked hydrogels (hereinafter abbreviated as ORgels) have many limitations in morphology and properties, e.g., morphological inhomogeneity, mechanicalweakness, limited swelling at equilibrium, and slow response to stimuli. Therefore, for example, NIPA-ORgels consisting of PNIPA chemical networks readily become turbid due to structural inhomogeneities inducedby increasing the crosslink density, pressure, and polymerization temperature [100–102]. Also, swelling ratiosor de-swelling rates are not always sufficient for certain applications [88, 103, 104]. These limitations mainly




exfoliatedclay

Dic

100 nm

g2

g1

χ

Figure 7.46 Schematic representation of the structural model with organic/inorganic networks in the NC gel.Dic is interparticle distance of exfoliated clay sheets; χ , g1 and g2 represent crosslinked chain, grafted chain, andlooped chain. Reprinted from [107]. Copyright (2002) with permission from John Wiley & Sons.

arise from restricted molecular motions of PNIPA chains caused by crosslinking with a large number oforganic crosslinkers. Furthermore, the most serious limitations of OR gels are due to their weak and brittlenature [105, 106]. Irrespective of the composition or conditions of preparation, NIPA-OR gels always broke atelongations less than 30% or when bent through 180◦ [107] and so were very difficult to handle in applicationsrequiring significant applied stress or strain. This limitation is due to the low average and broad distributionof chain lengths between crosslinking points in OR gels.

Recently, and in order to overcome the limitations of OR gels, a new type of polymer hydrogel nanocom-posite type hydrogels has been prepared (herein after abbreviated as NC gels, e.g., NIPA-NC gel) con-sisting of a unique organic polymer (e.g. PNIPA)/inorganic (clay) network (Figure 7.46) [107]. NC gelsshowed remarkable improvements in mechanical, optical, and swelling-deswelling properties. That is, NCgels simultaneously exhibit high transparency (structural homogeneity) irrespective of the preparation con-ditions, excellent tough mechanical properties, with astonishingly large elongations, large swelling ratios,and rapid deswelling responses to temperature changes. NIPA-NC gels were prepared by in situ free-radical




polymerization of N-isopropylacrylamide, in the presence of inorganic clay exfoliated in an aqueous media.The formation of effective polymer/clay networks could be achieved using inorganic clays which act as mul-tifunctional crosslinking agents through noncovalent interactions instead of using organic crosslinkers. Thecharacteristics of NC gels strongly depend on their compositions. In other words, they could be controlledover a wide range by altering their composition and network structure.

Haraguchi et al. prepared nanocomposite hydrogels composed of poly(N,N-dimethylacrylamide)(PDMAA) and clay by in situ free-radical polymerization of N,N-dimethylacrylamide (DMAA) in thepresence of inorganic clay in aqueous solution [108]. DMAA·NC gels consist of organic/inorganic networkssimilar to those of NIPA-NC gels established previously [107]. The proposed model structure shown inFigure 7.46 consists of networks of inorganic clay crosslinked by PDMAA through noncovalent interactions(probably hydrogen bonding ionic and/or coordinate interactions) between the PDMAA chains and the clay.Here, it is noted that, in the NC gels, free linear PDMAA chains could not be detected during the purificationand swelling processes. This suggests that all polymer chains are attached to the clay and involved in thenetwork. Also, the PDMAA component of the NC gel forms not only crosslinked chains (x) but also graftedchains with a free chain end (g1) or looped chains (g2) in which both ends are attached to one clay sheet.Further, as for crosslinked chains, topologically crosslinked chains consisting of looped chains or trappedinterchain entanglements, etc., might be included in NC gels, although the precise nature of the interactionsand the mechanism of network formation are still under examination. The interactions between (P)DMAAinitiator, and clay in the aqueous media and in the dried state were observed by similar means to those usedfor the PNIPA-c1ay system: viz. by solution viscometry and by IR spectroscopy. The proposed model for thenature of PDMAA/clay networks seems reasonable on the basis of analyses of their characteristic rubberymechanical properties and their swelling behavior.

To verify the proposed model, the clay dispersion and the chain flexibility of the polymer in DMAA-NCgels were examined. For the clay, the fact that the resulting NC gel and its dried gel were both opticallytransparent, almost regardless of clay content (Cclay), indicates that it is finely and uniformly dispersed in thehydrogels. The nature of the clay dispersion was also elucidated by SAXS measurements.

It was generally observed that DMAA-NC gels exhibited excellent tensile mechanical properties comparedwith conventional DMAA-OR gels. Similar to NIPA-OR gels, DMAA-OR gels broke at very low extension,ca. 50%, irrespective of crosslinker content (CBIS = 1 − 9) and showed very low ultimate tensile strengths(ca.7kPa). On the contrary, DMAA-NC gels generally exhibited very large elongations at break and highstrengths.

Also, it was observed that the tensile properties of DMAA-NC gels strongly depend on Cclay. Figure 7.47shows tensile stress–strain curves measured for DMAA-NC gels with different Cclay, where both modulus andstrength exhibit remarkable increases with increasing Cclay. On the other hand, elongations at break decreasedslightly with increasing Cclay, particularly at relatively low Cclay (< NC4). In samples with higher Cclay

(≥ NC4), the elongation at break changed little, staying at an almost constant value of ca. 1300%.The effects of clay content (Cclay) on the mechanical properties of DMAA-NC gels described above were

very similar to those for NIPA-NC gels [109], although the absolute values of properties and their detaileddependencies on composition were somewhat different. Changes in tensile modulus, tensile strength, andelongation at break by altering Cclay for both DMAA-NC-Ml and NIPA-NC-MI gels are summarized inFigure 7.48a–c. It was shown that DMAA-NC gels exhibit, as a whole, lower moduli and higher elongationsat break than those of NIPA-NC gels, while the ultimate strengths were similar for both gels.

The stress–strain behavior for a series of DMAA-NC2.5 and –NC5.5 gels with different polymer content(Cp), respectively were studied. Figure 7.49 shows the stress–strain curves for a series of DMAA-NC2.5 withdifferent (Cp). Cp was varied over about 2 orders of magnitude, from Cp = 0.1 to 8. NC gels prepared usingfixed conditions, apart from Cp, showed a large change in their mechanical characteristics, from very brittle tovery tough, on increasing Cp. In the very low Cp region, a lower critical polymer content (C*

p ), below which




250

200

150

NC7

NC4NC3

NC5.5

NC2.5

100

50

05000 1000

Strain (%)

Str

ess

[kP

a]

1500 2000

Figure 7.47 Stress–strain curves of DMAA-NC-Ml gels with different clay content (NC2.S to NC7). All hydrogelshad the same polymer/water ratio (= 1/10 (w/w). Reprinted from [108]. Copyright (2003) with permission fromAmerican Chemical Society.

NC gels became brittle, was observed at ca. 0.13 for both DMAA-NC2.5 and –NC5.5 gels. Hydrogels withCp < C*

p were too brittle to carry out standard tensile testing. Also, hydrogels prepared with Cp around C*p

were often neither uniform nor transparent. In the relatively low Cp region (C*p ∼ ca.0.5), the mechanical

properties of NC gels changed dramatically with increasing Cp. In particular, elongations at break increasedmarkedly from near zero to 1100–1600%. This indicates that the formation of the fundamental network ofNC gels requires this magnitude of Cp. Then, in the following intermediate Cp region, it was estimated thatincreasing Cp may result in an increase in the number of effective crosslinks. Actually, it was observed that themodulus and the strength clearly increased significantly with increasing Cp in this region (Cp = 0.5 – 5.5). Onthe other hand, the elongation at break was almost constant or changed little on increasing Cp. For example,in a series of NC2.5 gels, the elongations at break were observed to be constant at ca. 1600%, irrespective ofCp. However, in a series of NC5.5 gels, elongations at break gradually changed Cp in this region. Also, theabsolute maximum elongations of NC5.5 gels were, on the whole, a little less than those of NC2.5 gels.

Mechanical strength of polymer gel electrolytes was investigated on the basis of polymer-solvent affin-ity [110]. Four different polymers – poly(vinylidene fluoride) (PVdF), poly(vinylidene fluoride) (PVdF)-hexafluoropropylene (HFP) copolymer, poly(acrylonitrile) (PAN) and poly(methyl methacrylate) (PMMA) –were employed as the gel-forming polymer matrix, while ethylene carbonate (EC):propylene carbonate (PC)was used as the plasticizing solvent. The solvent retention ability of polymer gels decreases in the order ofPMMA ≥ PAN � P (VdF-HFP) ≥PVdF, which is surely the relative order of polymer affinity for the solvent.In Figure 7.50, the elastic modulus (slope of stress/strain curve in the elastic region), ultimate tensile strength(stress at break) and ultimate elongation (elongation at break) were compared between the polymer gel films.P(VdF-HFP) gel shows a higher elastic modulus and tensile strength as compared to the PMMA- or PAN-basedgels. It can thus be generalized that polymer gels are mechanically stronger when polymer matrix has a lowaffinity for solvent molecules, in which circumstances a microscopic phase transition is frequently observed.

7.5 Polymer Composites

During the last three decades, composites have replaced traditional materials in many engineering applicationsdue to their excellent specific properties. Fillers play important roles in modifying the desirable properties of




35

(a) Modulus

Mod

ulus

[kP

a]

30

25

20

15

10

5

0

(b) Strength

(c) Elongation at break

Str

engt

h [k

Pa]

Elo

ngat

ion

[%]

400

350

300

250

200

150

100

50

2000

1500

1000

500

0

2500

20 4

Cclay

6 8

0

DMAA-NC-M1 DMAA-NC4-M1* NIPA-NC-M1

Figure 7.48 Changes of (a) tensile modulus, (b) tensile strength, and (c) elongation at break by altering claycontents for two series of NC gels: DMAA-NC-MI gels (closed circle: solid line) and NIPA-NC-MI gels (opensquare: dotted line). The properties of DMAA-NC4-MI* gel with the same polymer weight with NIPA-NC4-MIgels are also plotted. Reprinted from [109]. Copyright (2003) with permission from American Chemical Society.

polymers and reducing the cost of their composites. The fibers such as glass fibers, aramide fibers, carbonfibers and natural fibers have been used as macrofillers to develop thermoplastic and thermoset composites.Inorganic fillers with dimensions in the micrometer range, e.g., calcium carbonate, glass beads, silica, carbonblack and talc, have been used extensively to enhance the mechanical properties of polymers. Such propertiescan indeed be tailored by changing the volume fraction, shape, and size of the filler particles. It is logical toanticipate that the dispersion of fillers with dimensions in the nanometer level having very large aspect ratio




250DMAA – NC2.5 gel

200

150

Cp = 0.13Cp = 0.22

Cp = 0.5

Cp = 4.5

Cp = 5.5Cp = 8

Cp = 1

Cp = 2

100

50

05000 1000

Strain (%)

Str

ess

[kP

a]

1500 2000

Figure 7.49 Effect of polymer contents (Cp = 0.13 − 8) on the stress–strain curves of DMAA-NC2.5 gels.All hydrogels had the same clay/water ratio (= 0.23/10 (w/w). Reprinted from [109]. Copyright (2003) withpermission from American Chemical Society.

and stiffness in a polymer matrix could lead to even higher mechanical performances. These fillers includelayered silicates, nanosilica and carbon nanotubes.

7.5.1 Mechanical properties of polymer macrocomposites

In the field of composites, the fiber reinforcement of matrices was initially developed using man-made fiberssuch as glass, carbon, aramid, etc., in order to take advantage of their high tensile moduli. Polymer-matrixcomposites, such as carbon or glass-fiber reinforced plastics (CFRP/GFRP) have been widely used in industrysince they have high strength and modulus. In fact, the total amount of consumption of GFRP was about 382thousand tons in Japan in 2001.

2.0

1.5

1.0

0.5

0.01000 200

Elongation/%

PAN gel PMMA gel

P(VdF-HFP) gel

Str

ess/

MP

a

300

Figure 7.50 Stress–strain curves of the polymer gel films. Reprinted from [110]. Copyright (2001) with permissionfrom Elsevier.




45

40

35

30

25

20

150 5 10 15

Cooling rate (°C/min)

SB

SS

(M

Pa)

20 25 30 35 40

m-PP/s-G

PP/G

Figure 7.51 Measured short beam shear stresses (SBSS) on the PP/G and m-PP/s-G composites as a function ofCR. Reprinted from [111]. Copyright (1998) with permission from John Wiley & Sons.

Youssef and Denault investigated the effect of the interfacial improvement by the addition of a chemicallymodified PP and a specific glass fiber thermoplastic sizing on the mechanical properties of the composites[111]. Mechanical tests were carried out on both PP/G (consisting of a pure PP matrix reinforced with 52 wt%of E-glass fibers) and m-PP/s-G (consisting of a blend of chemically modified PP with pure PP reinforcedwith 60 wt% of E-glass fibers coated with a specific thermoplastic sizing) composites. The study of theinterface quality via the measurement of the short beam stresses and the response of the composite undertension in the ±45◦ direction showed that the presence of the modified PP in the pure PP matrix and the fibersizing have noticeable effects on the interfacial strength. In fact, the m-PP/s-G system showed higher SBS(Short Beam Shear) stress values (Figure 7.51), higher tensile moduli and higher failure stresses in the ±45◦

direction. However, these performance parameters are significantly reduced when the composites are cooleddown very slowly resulting in large spherulitical crystalline structures. This suggests that the contribution ofthe amorphous phase to the fiber-matrix interaction in this composites is important.

When the thermosetting resins are used as a matrix for glass fiber, it is usually difficult to recycle thematerial. The thermoplastics are also impractical to recycle due to high cost and low quality issues in thepresent state. Most wasted FRPs are dumped although they do not decompose itself naturally in the ground,while others are burned. In recent years, a serious problem has come up in the use of plastics, especiallyfor the polymer composite materials. Energy recycling systems are also under development using polymercomposites as solid fuel. However, the glass fibers in the composites reduce the net heat and might damage thefurnace, if their composites are burned in it as a solid fuel. It introduces another problems such as the disposalof the remains, because the glass fibers in the composite would also remain in the incinerator. Meanwhile,natural fiber composites are claimed to offer environmental advantages such as reduced dependence on non-renewable energy/material sources, lower pollutant emissions, lower greenhouse gas emissions, enhancedenergy recovery, and end-of-life biodegradability of components. Such superior environmental performanceis an important driver of increased future use of natural fiber composites,

Tensile, flexural, and impact behavior of PALF-reinforced polyester composites as a function of fiberloading, fiber length, and fiber surface modification were investigated by Devi et al. [112]. The PALF–polyester




90FLEXURAL STRENGTH

FLEXURAL MODULUS

FLEXURAL STRENGTH FLEXURAL MODULUS

10 mm 30 mm

80

70

60

50

5

4

3

2

10

(a) (b)

10 20

FIBRE LENGTH, mm

FLE

XU

RA

L S

TR

EN

GT

H, M

Pa

30 40 50

100

60

80

40

20

0

5

4

3

2

10 5 10 15 20 25 30 35 40 45

FIBRE LOADING (wt %)F

LEX

UR

AL

ST

RE

NG

TH

, (M

Pa)

FLE

XU

RA

L M

OD

ULU

S, G

Pa

FLE

XU

RA

L M

OD

ULU

S (

GP

a)

Figure 7.52 (a) Variation of flexural strength and flexural modulus with fiber length (fiber loading 30 wt%); (b)Variation of flexural strength and flexural modulus with fiber loading of PALF–polyester composites [112].

composites exhibit superior mechanical properties when compared to other natural-fiber polyester compositesand can be used as structural composites. The optimum length of the fiber required to obtain PALF–polyestercomposites of maximum properties was found to be 30 mm (Figure 7.52a). The stress–strain behavior intension reveals that neat polyester is brittle and the addition of fibers makes the matrix more ductile. Thetensile strength and Young’s modulus of PALF polyester composites increased linearly with the fiber weightfraction. But in the case of flexural strength, there is a leveling off beyond 30% (Figure 7.52b). The impactstrength also increased linearly with the weight fraction of the fiber. The composite with 30 wt% fiber contentexhibits an impact strength of 24 kJ/m2. The high toughness of this natural fiber polymer composite places itin the category of tough engineering materials. A significant increase in the strength of the composites wasobserved after treatment of the fibers. The best improvement was observed in the case of silane A-172-treatedfiber composites.

Okubo et al. developed composites for ecological purposes (Eco-composites) using bamboo fibers extractedby the steam explosion technique [113]. The bamboo fibers (bundles) had a sufficient specific strength, whichis equivalent to that of conventional glass fibers. The tensile strength and modulus of PP-based compositesusing steam-exploded fibers increased about 15 and 30%, respectively, due to well impregnation and thereduction of the number of voids, compared to the composite using fibers that are mechanically extracted.The steam explosion technique is found to be an effective method to extract bamboo fibers for reinforcingthermoplastics

The influence of various chemical treatments on the properties of sisal/PE composites has been investigatedby Joseph et al. [114]. The chemical treatments included treatments with sodium hydroxide, isocyanate, andperoxide. The enhancement in the properties was ascribed to the bonding between sisal fiber and the PEmatrix. Treatment with the cardanol derivative of toluene isocyanate was found to be better than othertreatments as evidenced by the decrease in the hydrophilic nature of the composite. The composites exhibitedbetter dimensional stability and retention of properties even after aging, which was ascribed to the improvedmoisture resistance.




7.5.2 Mechanical Properties of Polymer Microcomposites

There is now considerable evidence that microfillers can significantly affect the structure of the matrix polymeritself, and hence the properties of the final composite. Commonly-used micro -fillers for polymers includemineral fillers such as talc, calcium carbonate, crystalline silica, and synthetic fillers such as carbon black,synthetic silica, etc.

The toughening of polypropylene with rigid particles leads to a system with higher stiffness and higherimpact resistance. A polypropylene–CaCO3 composite was processed by Zuiderduin et al. [115] which had asignificant higher modulus and simultaneously showed improved toughness. The notched Izod impact energycould be raised from 2 to 50– 60 kJ/m2 at room temperature while increasing the modulus.

Effect of processing conditions on the dynamic mechanical behavior of carbon black (CB) filled ethy-lene/ethylacrylate copolymer (EEA) composites was investigated [116]. The compounds were prepared bytwo methods, solution blending and mechanical mixing. Compared with the solution counterpart, the me-chanical composites have a high dynamic elastic modulus which results from the good dispersion state ofcarbon black in EEA, i.e., the strong interaction between carbon black and EEA.

7.5.3 Mechanical Properties of Polymer Nanocomposites

Polymeric nanocomposites can be considered as an important category of organic–inorganic hybrid materialsin which inorganic nanoscale building blocks (e.g., nanoparticles, nanotubes, or nanometer-thick sheets) aredispersed in an organic polymer matrix [117–121]. When compared to conventional composites based onmicrometer-sized fillers, the interface between the filler particles and the matrix in polymer nanocompositesconstitutes a much greater area within the bulk material, and hence influences the composite’s propertiesto a much greater extent, even at a rather low filler loading [122–124]. Polymer nanocomposites reinforcedby relatively small amounts of ultra-fine, nano-particles (most often clay platelets) proved exceptionallypromising engineering materials with an unexpectedly high stiffness/toughness ratio, gas barrier properties,flame retardence, etc. The real interest in nanotechnology is to create revolutionary properties and functionsby tailoring materials and designing devices on the nanometer scale.

According to a report, the total worldwide market for polymer nanocomposites reached 11.1 million Kgvalued to US$90.8 million in 2003. This market was expected to grow at an average annual growth rate of18.4% to reach US$211 million by 2008.

7.5.3.1 Factors Affecting Mechanical Properties of Nanocomposites

Influence of Filler Dispersion on Mechanical Properties Mechanical properties of polymer–claynanocomposites are highly related to their microstructure which in turn is directly related to the exfolia-tion and dispersion of clay platelets in the polymer matrix. Tensile properties of the MA compatibilizedPE–organoMMT nanocomposites were studied by Lee et al [125]. The tensile strength and modulus tend toincrease with increasing clay content (Figure 7.53a). Such an increasing trend is more obvious for the tensilemodulus. The increase in the tensile strength is higher at low clay content, indicating that the clay layers arebetter exfoliated. The reinforcing effect is lower for nanocomposites with higher clay content owing to someclay platelets being partially exfoliated and stacked. From Figure 7.53b, the strain at break decreases withincreasing clay content as expected.

The complete exfoliated morphology of polymer-MMT nanocomposites contributes greatly to their impactproperty and modulus, while the intercalated state of partially exfoliated state contributes more to the finalmaterial’s modulus. In the case of PP-based clay nanocomposite, how the dispersion state of MMTs decidesthe final properties of izod impact properties is shown in Figure 7.54 [126].




20

Tens

ile S

tren

gth

(MP

a)

Tensile Modulus (M

Pa)

15

0(a)

10

25

1 2 3Clay Content (%)

Tensile strengthTensile modulus

4 5 6 7 8

800

900

700

600

500

400

1000

(b)

150

Str

ain

at B

reak

(%

)

100

50

00

200

1 2 3Clay Content (%)

4 5 6 7 8

Figure 7.53 (a) Tensile modulus and strength and (b) strain at break vs. clay content for the MAcompatibilizedPE–organoMMT nanocomposites. Reprinted from [125]. Copyright (2005) with permission from Elsevier.

Liu et al. synthesized organoclay-modified epoxy nanocomposites by high pressure mixing method(HPMM) [116]. Better dispersion of organoclay in epoxy matrix has been achieved by HPMM than di-rect mixing method (DMM) (Figure 7.55). The nanocomposites formed by the HPMM showed a dramaticimprovement in fracture toughness at very low clay loading; that is, K1C and G1C were increased by 1.7 and3.2 times respectively, at 1.5 phr (about 1 wt%) organoclay loading (Figure 7.56).

44

40

Exfoliated state

Blending Intercalated state

36

32

Izod

Impa

ct S

tren

gth

(J/m

)

28

242 4

Clay Content (%)6 8 100

Figure 7.54 Relationship between the izod impact property–clay load in polypropylene-based nanocomposites.Reprinted from [126]. Copyright (2005) Elsevier.




SEI 15.0 kV x550 20μm SEI 15.0 kV x2000 5μm

Figure 7.55 SEM micrographs of fracture surfaces of nanocomposites made with (a) DMM showing agglomeratesclay layers in epoxy matrix; (b) HPMM showing fine dispersion of clay layers in the epoxy matrix. Reprinted from[127]. Copyright (2005) with permission from Elsevier.

Influence of Filler Content on the Mechanical Properties The amount of filler strongly influences themechanical properties of polymer nanocomposites. Zhou et al. studied the effect of filler content on themechanical properties of multi-walled carbon nanotubes (CNTs) reinforced epoxy [128]. When dispersingCNT in polymer matrix, it is important to keep the filler volume (or weight) fraction below a certain valueto maintain the strength and fracture toughness. Optimal loading of CNT in matrix is a key parameter todeveloping multifunctional nanophased composite. Flexural modulus steadily increases with a higher CNTweight percent. Modulus improved by 11.7% with an addition of a 0.4 wt% of CNTs. Flexural strength andfracture toughness peaked in a 0.3 wt% CNT/epoxy system. The decrease in strength and fracture toughness in0.4% CNT/epoxy was attributed to poor dispersions of nanotubes in the composite. Compared to neat epoxy,

5001.30E with HPMM

1.30E with DMM

Na-Mont400

300

200

100

00 4

Clay Loading, phr

(a) (b)

G1C

, J/m

2

8 12 16

1.30E with HPMM

1.30E with DMM

Na-Mont

0.5

0.75

1

1.25

1.5

1.75

2

0.250 4

Clay Loading, phr

K1C

, Mpa

. m1/

2

8 12 16

Figure 7.56 (a) Critical strain energy release rate (G1C); (b) Critical stress intensity factor (K1C) of nanocompos-ites and filler composites [127].




Table 7.6 Mechanical properties of neat and CNT modified epoxy [128].

Modulus [GPa] Strength [MPa] Failure strain [%]

Neat epoxy 2.46 93.5 4.020.1% CNT 2.54 109 6.060.2% CNT 2.60 115 6.800.3% CNT 2.65 121 7.580.4% CNT 2.75 113 5.12

DMA results indicated a 93% improvement in storage modulus in 0.4 wt% CNT/epoxy at room temperatureand a 17◦C increase in Tg (Table 7.6, Figure 7.57).

The approach for incorporating nanoscopic inorganic cluster into organic polymer is to design well-definedinorganic oligomers with a single polymerization site per cluster. Each oligomeric cluster has an exactly de-fined degree of polymerization of eight, (RSiO1.5)8, or more precisely, P1R7Si8O12, where R and P are organicgroups. These polyhedral oligomeric silsesquioxane (POSS) macromers have an inorganic silica-like coreand are surrounded by eight organic groups, of which seven are inert and just one is reactive. Further poly-merization involving the single reactive P site results in a linear polymer containing monodisperse, nanosizeinorganic clusters pendent to an organic polymer backbone. These hybrid inorganic–organic polymers can beprocessed further like any thermoplastic polymers. Lee and Lichtenhan have modified epoxy with polymer-izable polyhedral oligomeric silsesquioxane (POSS) macromers [129]. Figure 7.58, depicts the relaxationmodulus curves tested for different loading of POSS–epoxy monomers after 64 h of isothermal aging at atemperature of 63.9◦C. It is clear that these relaxation curves can be superimposed with only horizontal shiftsalong the time axis. However, it is interesting to point out that the value of E0 is not affected by the pres-ence of the nanoreinforcement. This may be in part a reflection of the monofunctional nature in POSS epoxy

180

160

140

Frac

ture

toug

hnes

s [M

Pa.

mm

1/2 ]

120

1000.00 0.10 0.20

Weight fraction of CNTs [%]0.30 0.40

Figure 7.57 Effect of CNT contents on fracture toughness of epoxy. Reprinted from [128]. Copyright (2008) withpermission from eXPRESS Polymer Letters.




100

101DGEBA/D230Heloxy67x%POSS

10–2

10–1 100 101 102

Time (Seconds)

103 104 105

10–1

0%POSS

5%POSS

9%POSS

ε=0.1%

T=63.9°C

ta=64 hours

Rel

axat

ion

Mod

ulus

(G

Pa)

Figure 7.58 Small-strain stress relaxation modulus curves for DGEBA–D230–Heloxy 67–wt % POSS–epoxy glassafter 64 h of isothermal aging at a temperature of 63.9◦C. Applied strain is 0.001[129].

monomers used in this study. Despite the ability of the POSS cages to hinder the relaxation motion of networkjunctions from a chain terminus location within the network, they do not contribute to the overall deformationprocess of such glassy networks from this position. Interestingly, such monofunctional POSS–epoxide maybe useful for enhancing glass transition without increasing crosslink density and potentially detracting fromthe desirable mechanical properties of such epoxy networks.

Filler-Matrix Adhesion The effect of clay modification on organo-MMT/NBR nanocomposites was stud-ied by Kim et al. [130]. They modified the organoclays with alkylamine cations. Alkylamines used wereoctylamine (C8), DDA (C12), and ODA (C18). The organo-MMT content in the prepared nanocompositeswas fixed at 0.0, 1.86, 4.52, 8.70, 12.45, and 15.94 wt%. For the C8-MMT/NBR nanocomposites, the tensilestrength increased rapidly with increasing clay content from 0 to 4.52 wt%, but the change was less whenthe clay content increased beyond 4.52 wt%. In the case of C12-MMT and C18-MMT, the tensile strengthincreased rapidly with the clay content up to 8.7 wt% beyond which there was very little change. For theC8-MMT/NBR system, the tensile modulus increased slightly with increasing clay content. But the tensilemodulus of C12-MMT/NBR and C18-MMT/NBR nanocomposites increased rapidly with increasing claycontent. Below 8.7 wt%, the improvement in elongation at break may be attributed in part to the plasticizingeffect of alkyl ammonium ions that are located at the clay–NBR interface. Above 8.7 wt%, the lowering ofthe elongation at break was explained by the formation of nonexfoliated aggregates at higher clay content thatmade these composites much stiffer. The differences in mechanical properties among the C8-MMT/NBR,C12-MMT/NBR, and C18-MMT/NBR hybrids were explained by the differences in hydrophobicity of the




135130125120115110105100

0 2 4clay content (wt%)

Fle

xura

l str

engt

h (M

pa)

6 8 10

E+OC seriesE+UC series

4.5

4

3.5

3

2.5

20 2 4

clay content (wt%)

Fle

xure

mod

ulus

(G

Pa)

6 8 10

E+OC seriesE+UC series

Figure 7.59 (a) Flexural strength of epoxy/clay series; (b) Flexural modulus of epoxy/clay series [131].

organo-MMT. Overall, the mechanical properties increased in the order C8-MMT < C12-MMT < C18-MMT,depending on the length of the alkyl chain in the alkyl ammonium.

The diglycidyl ether of bisphenol A (DGEBA) epoxy resin system filled with organo clay (OC) andunmodified clay (UC) were processed and Mohan et al. studied the effect of filler-matrix adhesion onmechanical properties [131]. The flexural properties of the epoxy filled with organo and unmodified clayparticles are shown in Figure 7.59.

The organoclay-filled epoxy has better improvement in the flexural modulus than that of the unmodifiedclay-filled epoxy. The important parameter that affects this property by incorporating such fillers is thequality of interface in the composites, i.e., the adhesive strength and the interfacial stiffness of the compositemedium. These two factors play a crucial role in stress transfer and the elastic deformation from the matrixto the fillers. This is very much applicable to the nanoparticle-filled polymers, due to high surface area ofthe organoclay filler which increases the contact area to the matrix and imparts a high portion of interface.If the interface is poor between the matrix and fillers, there is much less chance for the fillers to carry theload and results in less modulus, which is seen for the unmodified clay-filled epoxy polymer. As a resultof this, the unmodified clay-filled composite cannot have good strength as compared to the matrix material.But in the nanocomposites, because of the enhanced interfacial property owing to their large surface area offillers, the increased property is observed which reveals that stresses are effectively transferred through theinterface.

For polymer-based nanocomposites, an appropriate surface treatment of inorganic nanoparticles shouldnot only improve dispersion of the fillers, but also bring about notable influence on the interfacial character-istics, and subsequently enhance the mechanical properties of the ultimate composites. The internal surfaces(interfaces) are critical in determining the properties of polymer silica nanocomposites. Taking the advantageof graft polymerization, Cai et al. [132] prepared nanosilica-modified polypropylene with specific structuralrequirements in one step. The key issue lies in the introduction of a polymeric foaming agent containingsoft segments (i.e., poly(p-vinylphenylsulfonylhydrazide-co-butyl acrylate)) onto the surface of nano-SiO2.In the case of melt blending with PP, the side sulfonyl hydrazide groups on the grafted copolymer are ableto be gasified like foaming agent to induce a localized bubble-stretching effect that pulls apart nanoparticlesagglomerates. Meanwhile, the skeleton of the grafted copolymer would get entangled with the matrix polymerforming strong interfacial interaction, and the poly (butyl acrylate) units in the grafted copolymer help toraise ductility of the interlayer (Figure 7.60).

Compared to the composite containing rigid macromolecular foaming agent grafted nano-SiO2 (i.e., poly(p-vinylphenylsulfonylhydrazide) grafted nano-SiO2/PP), the composite fabricated in this work (i.e., poly(pvinylphenylsulfonylhydrazide-co-butyl acrylate) grafted nano-SiO2/PP) shows much greater increment innotched impact strength without expense of lowing tensile performance.




Residual sulfonyl hydrazide group

PBA chain unitsO

O

O

mn

CH2

SO2NHNH2

OCH2 OCH2Si(C

H2 )3

CH2

CH

CH

C

C C4H9

O = O

PP matrixOH

Graft copolymerization

In situ bubble stretching

Bubble-stretching induced dispersionof SiO2 nanoparticles in PP matrix andsimultaneous interfacial tailoring

SiO2 nanoparticles Bubble generated by gasificationof side sulfonyl hydrazide groupson the grafted polymer

Figure 7.60 Schematic drawing of the proposed route for making nano-SiO2/PP composites [60].

7.5.4 Mechanical Modeling of Polymer Nanocomposites

Several models have been developed to predict the mechanical behavior of polymer nanocomposites, includingHalpin–Tsai, Mori–Tanaka, etc. The models generally include parameters such as the aspect ratio, volumefraction and the orientation of the reinforcement.

The elastic modulus of the composite material can be predicted from the Halpin–Tsai equation assumingthe fibers are discontinuous and aligned uniaxially [133, 134]. The longitudinal elastic modulus of compositesEc, is given by:

Ec

Em= 1 + ξηφ

1 − ηφ(7.15)

where Em is the tensile modulus of the matrix and ϕ is the volume fraction of fiber reinforcement.The Mori–Tanaka mean field theory is used to assess the overall properties such as the effective stiffness

tensor C* of the composites. It is based on the Elsheby method for estimating stress state in compositereinforced with misfitting inclusions. The composite is assumed to be composed of a continuous matrix anddiscrete of inclusions of different stiffness. The effective stiffness tensor C* is given by the following relation[135, 136]:

C∗ = C1 + V2{(C2 − C1)} (7.16)

where C1 is the matrix phase stiffness tensor, C2 the inclusion stiffness tensor, V2 the inclusion volume ratio,and A is the concentration tensor.

For a composite consisting of a single, arbitrarily shaped inclusion perfectly bonded inside the matrix, thedilute strain concentration tensor of the effective particle is given by:

A(di1) = [I + SC−1(C2 − C1)]−1 (7.17)




where I is the fourth order unit tensor and S the fourth order Elshelby tensor. As the inclusion volume fractionincreases, interaction between the inclusions reduces the accuracy of the dilute approximation.In other words,interactions of the field from other inclusions are expected to influence the evolution of the average fieldsin the matrix and the reinforcement. The Mori–Tanaka approach includes the effect of particle interaction.[137] In this case, A can be expressed as:

A = A(di1)[V1 I + V2{A(di1)}]−1 (7.18)

where V1 is the matrix volume ratio. The Mori–Tanaka model has better predictive capability for fillerswith relatively high aspect ratios. Tandon and Weng based on the Mori–Tanaka approach and derived thelongitudinal modulus (E11) of the composite reinforced with platelets [138]:

E11

Em= 1

1 + φ f [−2νm A3 + (1 − νm)A4 + (1 + νm)A5 A]/2A(7.19)

where νm is the Poisson’s ratio of the matrix, and A, A3, A4 and A5 are calculated from the matrix and fillerproperties and the components of the Elshelby tensor.

In polymer–clay nanocomposites, parameters associated with hierarchical morphology of the clay such asthe silicate interlayer spacing (d0 0 1), gallery spacing, platelet thickness, etc., should be incorporated into themicromechanics model. Brune and Bicerano [139] modified the Halpin–Tsai equation for tensile modulus ofintercalated (or incompletely exfoliated) nanocomposites as:

Ec

Em= 1 + ξ ′η′φ′

1 − η′φ′ (7.20)

where Er′ is the ratio of the modulus of the platelet stack to that of the matrix, ξ ′ the aspect ratio of the platelet

stack and η′ is the volume fraction of platelet stacks in the matrix.Recently, Fornes and Paul [140] emphasized the importance of aspect ratio and exfoliation ratio of clay

in modeling stiffness. They performed simple calculations on the aspect ratio (l/t) of MMT platelets ofthe PA6–organoclay nanocomposites. They quantified those parameters and used the composite theories ofHalpin–Tsai and Mori–Tanaka to predict the stiffness of high molecular weight nylon 6/clay nanocompositesand obtained good agreement with experimental data (Figure 7.61). They assumed that the composite consistsof a matrix and stacks of clay platelet sheets. For the particle thickness (t) determination, they incorporatedseveral parameters such as the silicate interlayer spacing (d0 0 1), number of platelets per particle and thethickness of an MMT platelet. From this, the number average particle thickness was determined to be 1.61nm. On the basis of image analysis from TEM micrographs of nanocomposites, they determined the averageparticle length to be 91 nm. This resulted in an aspect ratio of 57. For an exfoliated structure, the plateletswere completely delaminated and dispersed independently in the matrix with a thickness of 0.94 nm. Thusan aspect ratio of 97 was determined for an exfoliated structure. They substituted such particle aspect ratiovalues together with the stiffness of MMT (178 GPa) and PA6 (2.75 GPa) into the Halpin–Tsai equation. InFigure 7.61(a), experimental modulus data was compared with model predictions for aligned layered alumino-silicate nanocomposites having an aspect ratio identical to the experimental value of 57 (number averageaspect ratio). The ratio of 97 corresponds to perfectly exfoliated morphology, i.e., the number average particlelength, 91 nm, divided by the thickness of an individual platelet, 0.94 nm. It is clear from Figure 7.61(a) thatHalpin–Tsai equations slightly overpredict the experimental data, while the Mori–Tanaka theory underpredictsthe experimental data. However, both theories demonstrate that even higher levels of reinforcement might bepossible with higher levels of exfoliation, larger platelet diameters, improved orientation, etc. Figures 7.61(b)and (c) compare the experimental data to model predictions based on an intercalated morphology (stacks of




3

2

1

E/E

m0 1 2 3 4

Halpin-Tsai

Mori-Tanaka

vol % MMT

1 / t97

57

97

57

HMW/(HE)2M1R1 Expt. Data

Ef = 178 GPaEm = 2.75 GPavf = 0.20vm = 0.35

(a)

3

2

1

E/E

m

0 1 2 3 4

Halpin-Tsai Equations

vol % MMT

1

23

n

HMW/(HE)2M3R3 Data (n=1.4)

Em = 2.75 GPat = 91 nmdm = 1.3 nm

(b)

3

2

1

E/E

m

0 1 2 3 4

vol % MMT

HMW/(HE)2M1R3 Data (n=1.4)

Mori-Tanaka TheoryEm = 2.75 GPat = 91 nmdm = 1.3 nm

1

23

n

(c)

Figure 7.61 Comparison of experimental and theoretical (based on unidirectional reinforcement) stiffness forhigh molecular weight nylon 6 nanocomposites. (a) Pure montmorillonite having a filler modulus of 178 GPaand aspect ratio of 57 (experimentally determined number average value) and 97, corresponding to completeexfoliation, and (b, c) intercalated morphology having one or more platelets per stack [129].

clay intercalated with polymer having one or more platelets per stack). The stack properties were based onexperimental data, i.e., the stacks are 91 nm in length, have a repeat spacing of 1.8 nm, and each individualdisk has a modulus of 178 GPa. It can be seen that the experimental data lie between the Halpin–Tsai curvescorresponding to 1 and 2 platelets per stack, which is very close to the experimental determined value of1.4. However, when Mori–Tanaka theory is used, the experimental data matches a completely exfoliatedmorphology, i.e., n = 1.




7.6 Conclusion, Future Trends and Challenges

The research in the area of multiphase polymer systems, especially polymer blends and composites, hasgot much attention for the past several decades in polymer science and technology. In this chapter, wehave presented the various aspects concerning the mechanical and viscoelastic properties of multiphasepolymer systems based on thermoplastics, rubbers and thermosets. The existing literature on mechanicalproperties of polymer blends including nanostuctured blends, polymer composites, IPNs and polymer gelshave been discussed. The critical parameters which tune the mechanical properties of multiphase system aredemonstrated with examples. The applicability of various micromechanical models on different multiphasepolymer systems is included in this chapter.

More effort needs to be placed in the field of computation modeling aiming at predicting the mechanical be-havior of multiphase polymer systems. Detailed studies on the mechanical modeling of polymer gels are alsorequired for predicting the mechanical properties. It should be emphasized that there is limited work on analyt-ical modeling of the strength and toughness of polymer nanocomposites. There are several conflicting reportson the mechanical properties of polymer nanocomposites. More careful and in-depth analysis should be doneto correlate the microstructure and mechanical properties of polymer nanocomposites. It is also important tostudy the effect of different mixers on the viscosity and hence on the morphology and properties of multiphasepolymer systems. For the better marketing of the products we also need to look at the cost–performance ratio.Finally, it is important to add that a comprehensive understanding of microstructure and mechanical propertiesis very important for the fabrication of useful products from multiphase polymer systems.

References

1. L. A. Utracki, in Polymer Blends Handbook, L.A. Utracki (Ed.), Volume 1, Kluwer Academic Publishers, Dordrecht(2002).

2. I. W. Hamley, The Physics of Block Copolymers, Oxford, Oxford Science Publications (1998).3. S. T. Milner, Macromolecules, 27, 2333 (1994).4. H. Bramfeldt, P. Sarazin and P. Vermette, Polymer Degradation and Stability, 93, 877–882 (2008).5. A. Kelarakis and K. Yoon, European Polymer Journal, 44, 3941–3945 (2008).6. P. Tanpaiboonkul, W. Lerdwijitjarud, A. Sirivat and R.G. Larson, Polymer, 48, 3822–3835 (2007).7. B. G. Sumpter, D. W. Noid and M. D. Barnes, Polymer, 44, 4389–4403 (2003).8. C. Harrats, S. Thomas and G. Groeninckx, Micro and Nanostructured Multiphase Polymer Blend Systems: Phase

Morphology and Interfaces, Taylor and Francis Group (2006).9. D. S. Lee and S. C. Kim, Macromolecules, 17, 11 (1984).

10. G. M. Kavanagh, S. B. Ross-Murphy, Prog. Polym. Sci., 23, 533–562 (1998).11. Y. Osada and J. P. Gong, Adv. Mater., 10, 11 (1998).12. S. G. Advani, Processing and Properties of Nanocomposites, World Scientific Publishing Co. Pte. Ltd (2007).13. D. R. Paul and C. B. Bucknall (Eds). Polymer Blends (Vol.1), Formulation (Vol.2), Performance, John Wiley &

Sons, Inc, New York (2000).14. C. W. Macoscko, Macromol.Symp., 149, 171–184 (2000).15. H. Varghese, S. S. Bhagawan, S. S. Rao and S.Thomas, Eur. Polym. J., 31, 10, 957–967 (1995).16. C. R. Kumar, K. E. George and S. Thomas, Journal of Applied Polymer Science, 61, 2383–2396 (1996).17. H. Veenstra, P. C. J. Verkooijen, B. J. J. van Lent, J. van Dam, A. P. de Boer and A. P. H. J. Nijhof, Polymer, 41,

1817–1826 (2000).18. Z. Oommen and S. Thomas, J. Appl. Polym. Sci., 65, 1245 (1997).19. M. E. Broz, D. L. Vander Hart and N. R. Washburn, Biomaterials, 24, 4181–4190 (2003).20. S. C. George, K. N. Ninan, G. Groeninckx and S. Thomas, Journal of Applied Polymer Science, 78, 1280–1303

(2000).




21. T. Semba, K. Kitagawa, U. S. Ishiaku and H. Hamada, Journal of Applied Polymer Science, 101, 1816–1825 (2006).

22. J. H. Chen, J. C. Zhong, Y. H. Cai, W. B. Su and Y. B. Yang, Polymer, 48, 2946–2957 (2007).23. L. A. Utracki and R. Simha, Macromolecules, 37, 10123 (2004).24. S. D. Burnside and E. P. Giannelis, Chem Mater., 7, 1596 (1995).25. C. Creton, E. J. Kramer, H. R. Brown and C. Y. Hui, Adv. Polym. Sci., 156, 53 (2001).26. S. Jose, PhD thesis, Mahatma Gandhi University (2006).27. J. Jaafar, A. F. Ismail and T. Matsuura, Journal of Membrane Science, 345, 119–127 (2009)28. G. Spadaro and G. Rizzo, Eur. Polym, J., 25, 1189 (1989).29. S. J. Kim, S. R. Chowduhury, W. J. Cho and C. S. Ha, J. Appl. Polym. Sci., 89, 5, 1305 (2003).30. N. Moussaif and G. Groeninckx, Polymer, 44, 7899 (2003).31. G. Filippone, N. T. Dintcheva, D. Acierno and F. P. La Mantia, Polymer, 49, 1312 (2008).32. M. Feng, F. L. Gong, C. Zhao, G. Chen, S. Zhang and M. Yang, Polym. Int., 53, 1529 (2004).33. Y.J. Li and H. Shimizu, Macromol. Rapid Commun., 26, 710 (2005).34. B. B. Khatua, D. J. Lee, H. Y. Kim and J. K. Kim, Macromolecules, 37, 2454 (2004).35. F. C. Chiu, H. Z. Yen and C. E. Lee, Polymer Testing, 29, 397–406 (2010).36. P. Motamedi and R. Bagher, Materials & Design, (2010).37. A. Rahmatpour, M. Abdollahi and M. Shojaee, Journal of Macromolecular Science, Part B: Physics, 47, 523–531

(2008).38. L. T. Vo and E. P. Giannelis, Macromolecules, 40, 8271–8276 (2007).39. S. Joseph and S. Thomas, J. Polym. Sci. B: Polym. Phys., 40, 755–764 (2002).40. R. C. Willemse, A. Speijer, A. E. Langeraar and A. Posthuma de Boer, Polymer, 40, 6645–6650 (1999).41. A. Y. Coran and R. Patel, J. Appl Polym. Sci., 20, 3005 (1976).42. T. Kunori and P. H. Geil, J Macromol Sci Phys B, 18, 135 (1980).43. E. H. Kerner, Proc. Phys Soc., 696, 808, (1956).44. A. H. J. Nijhof, LTM-report 1167, Delft University of Technology, The Netherlands, 1998.45. J. J. Rajasekaran, J. G. Curro and J. D. Heneycutt, Macromolecules, 28, 6843 (1995).46. M. Wang and N. Pan, Materials Science and Engineering R, 63, 1–30 (2008).47. B. D. Favis, J. Appl. Polym.Sci., 39, 285 (1990).48. J. Lu, Z. Qiu and W. Yang, Polymer, 48, 4196–4204 (2007).49. J. M. Raj and C. Ranganathaiah, Polymer Degradation and Stability, 94, 397–403 (2009).50. J. M. Huang and S. J. Yang, Polymer, 46, 8068–8078 (2005).51. A. Robard, D. Patterson and G. Delmas, Macromolecules, 10, 706 (1977).52. L. E. Fried, M. R. Manna and P. F. Pagoria, Annual Review of Materials Research, 31, 29, (2001).53. S. Joseph, Ph.D thesis, Mahatma Gandhi University (2004).54. K. A. Moly, S. S. Bhagawan, G. Groeninckx and S. Thomas, J. Appl. Polym. Sci., 100, 4526 (2006).55. B. John, K. T. Varughes, Z. Oommen, P. Potschke and S. Thomas, Journal of Applied Polymer Science, 87, 2083–2099

(2003).56. B. Baghaei, S. H. Jafari, H. A. Khonakdar, I. Rezaeian, L. As’habi and S. Ahmadian, Polym. Bull., 62, 255–270

(2009).57. B. K. Kim, Y. S. Oh, Y. M. Lee, L. K. Yoon and S. Lee, Polymer, 41, 385–390 (2000).58. J. Karger-Kocsis and V.N. Kuleznev, Polymer, 23, 699–705 (1982).59. R. J. Gaymans and J. W. van der Werff, Polymer, 35, 3658–3664 (1994).60. E. Biber, G. Gunduz, B. Mavis and U. Colak, Materials Chemistry and Physics, 122, 93–101 (2010).61. C. B. Bucknall, Toughened Plastics, Applied Science, London (1977).62. S. Wu, J. Appl. Polym.Sci., 26, 1855–1863 (1985).63. C. B. Bucknall and D. R. Paul, Polymer, 50, 5539–5548 (2009).64. R. J. M. Borggreve and R. J. Gaymans, Polymer, 30, 63–70 (1989).65. A.Yu. Vasilenko, D. D. Novikov and E. V. Prut, Doklady Physical Chemistry, 428, 2 (2009).66. A. Gonzalez-Montiel, H. Keskkula and D. R. Paul, Polymer, 36, 24, 4587–4603 (1995).67. G. Tripathi and D. Srivastava, Materials Science and Engineering A 443, 262–269 (2007).




68. R. Thomas, J. Abraham, P. Thomas, and S. Thomas, Journal of Polymer Science: Part B: Polymer Physics, 42,2531–2544 (2004).

69. S. Newman and S. Strella, J. Appl. Polym. Sci., 9, 2297 (1965).70. T. Kietzke, D. Nehe and K. Landfester, Novel approaches to polymer blends based on polymer nanoprticles, Nature

Mater., 2, 48–412 (2003).71. S. Ritzenthaler, F. Court, E. Girard-Reydet, L. Leibler and J. P. Pascault, Macromolecules, 36, 118–126 (2003).72. Y.S . Thio, J. Wu and F. S. Bates, Macromolecules, 39, 21, 7187–7189 (2006).73. L. H. Sperling, Thermodynamics and kinetics of phase separation, in Interpenetrating Polymer Networks, D.

Klemper, L. H. Sperling and L. A. Utracki (Eds), American Chemical Society, Washington, 77–123 (1994).74. L. A. Utracki, Thermodynamics and kinetics of phase separation, in Interpenetrating Polymer Networks, D. Klemper,

L. H. Sperling and L. A. Utracki (Eds), American Chemical Society, Washington, 77–123 (1994).75. J. Karger Kocsis, O. Gryshchuk. and S. Schimitt, J. Mater. Sci., 38, 413–420 (2003).76. D. Sopphiea, D. Klempner, V. Sendijarevic, B. Sutha, and K. C. Frisch, Interpenetrating polymer networks as

energy-absorbing materials, in Interpenetrating Polymer Networks, D. Klemper, L. H. Sperling and L.A. Utracki(Eds), American Chemical Society, Washington, 39–75 (1994).

77. D. Klempner and D. Sophiea, Interpenetrating polymer networks, in Elastomer Technology Handbooks, N. P.Cheremisinoff (Ed.), CRC Press, Boca Raton, FL, 421–444 (1993).

78. J. Karger-Kocsis, O. Gryshchuk and N. Jost., J. Appl. Polym. Sci., 88, 2124–2131 (2003).79. S. C. Kim, D. Klempner, K. C. Frisch and H. L. Frisch, Macromolecules, 10, 6 (1977).80. V. Huelck, D. A. Thomas and L. H. Sperling, Macromolecules, 5, 4 (1972).81. J. John, R. Suriyakala, S. Thomas, J. M. Mendez, A. Pius and S. Thomas, Journal of Materials Science (2010).82. J. John, D. Klepac, M. Didovi, C. J. Sandesh, Y. Liu, K. V. S. N. Raju, A. Pius, S. Valic and S. Thomas, Polymer,

51, 2390–2402 (2010).83. P. Bauer, J. Henning, and G. Schreyer, Angew. Makromol, Chem., 11, 145 (1970).84. D. De Rossi, K. Kajiwara, Y. Osada. and A. Yamauchi (Eds), Polymer Gels, Plenum: New York (1991).85. Y. Hirokawa and Y. Tanaka, J. Chern. Phys., 89, 6379 (1988).86. E. S. Matuo and T. J. Tanaka, Chern. Phys., 89, 1695 (1988).87. A. Suzuki and T. Tanaka, Nature (London), 346, 345 (1990).88. Y. H. Bae, T. Okano and S. W. Kim, J. Polym. Sci., Part B: Polym. Phys., 28, 923–936 (1990).89. K. Otake, H. Inomata, M. Konno and S. Saito, Macromolecules, 23, 283 (1990).90. T. Okano, Y. H. Bae, H. Jacobs and S. W. Kim, J. Controlled Release, II, 255–265 (1990).91. H. Inomata, S. Goto, K. Otake, and S. Saito, Langmuir, 8, 687–690 (1992).92. M. Shibayama, M. Morimoto and S. Nomura, Macromolecules, 27, 5060–5066 (1994).93. E. Kato, J. Chem. Phys., 106, 3792–3797 (1997).94. S. Hirotsu and A. Onuki, J. Phys. Soc. Jpn, 58, 1508–1511 (1989).95. F. Afroze, E. Niles and H. Berghamns, J.Mol.Struct., 554, 54 (2000).96. D. Dhara and P. R.Chaterji, Polymer, 41, 6133 (2000).97. M. Annkaka, K. Motokawa, S. Sasaki, T. Nakahira, H. Kawasaki, H. Maeda, Y. Amo and Y. Tominaga, J.Chem.

Phys., 113, 5980 (2000).98. Y. Li and T. Tanaka, J.Chem. Phys. 41, 6133 (1980).99. L. Liang, P. C. Roeke, J. Liu, G. E. Fryxell, J. S. Young, M. H. Engelhard and K. L. Alford, Langmuir, 16, 8016

(2000).100. T. Okajima, I. Harada, K. Nishio and S.Hirotsu, Jpn. J. Appl. Phy., 39, 1875 (2000).101. C. Nakamoto, T. Motonaga and M. Shibayama, Macromolecules, 34, 911 (2001).102. M. Shibayama, S. Tanaka and T. Norisuye, Physics A, 249, 245 (1998).103. R. Yoshida, K. Uchida, Y. Kaneko, K. Sakai, A. Kikuchi, Y. Sakuraia and T.Okano, Nature (London), 374, 249

(1995).104. Y. Kaneko, S. Nakamura, K. Sakai, A. Kikuchi, T. Aoyagi, Y. Sakkurai and T. Okano, J. Biomater. Sci. Polym. Ed.,

10, 1079 (1999).105. S. Hirotsu and A. Onuk, J. Phy. Soc. Jpn., 58, 1508 (1988).106. T. Takigawa, H. Araki, K. Takahashi and T. Masuda, J. Chem. Phys., 113, 7640 (2000).




107. K. Haraguchi and T. Takehisa, Adv. Mater., 14, 1120 (2002).108. K. Haraguchi, R. Farnworth, A. Ohbayashi and T. Takehisa, Macromolecules, 36, 15, 10162 (2003).109. K. Haraguchi, T. Takehisa, S. Fan, Macromolecules, 35, 10162–10171 (2002).110. C. S. Kim and S. M. Oh, Electrochimica Acta, 46, 1323–1331 (2001).111. Y. Youssef and J. Denault, Polymer Composites, 19, 3 (1988).112. L. U. Devi, S. S. Bhagawan and S. Thomas, J Appl Polym Sci., 64, 1739 ( 1997).113. K. Okubo,T. Fujii and Y. Yamamoto, Composites: Part A, 35, 377–383 (2004).114. K. Joseph, S. Thomas and C. Pavithran, Polymer, 37, 5139 (1996).115. W. C. J. Zuiderduin, C. Westzaan, J. Huetink and R. J. Gaymans, Polymer, 44, 261–275 (2003).116. H. Tang, X. Chen and Y. Luo, Eur. Polym. J., 32, 8, 963–966 (1996).117. G. Kickelbick, Prog. Polym. Sci., 28, 83 (2003).118. M. Yoshida, M. Lal, N. D. Kumar and P. N. Prasad, J. Mater. Sci., 32, 4047 (1997).119. T. J. Pinnavaia and G. W. Beall, Polymer-Clay Nanocomposites, John Wiley & Sons Ltd., Chichester, UK, (2001).120. K. Friedrich, S. Fakirov and Z. Zhang, (Eds), Polymer Composites – From Nano- to Macro Scale, Springer, NY,

USA (2005).121. G. Schmidt and M. M. Malwitz, Current Opinion in Colloid and Interface Science, 8, 103 (2003).122. W. Wang, Y. Li, L. Wei and Y. Fang, Mater. Lett., 57, 3366 (2003).123. S. Varghese, K. G. Gatose, A. A. Apostolov and J. Karger-Kocsis, J. Appl. Polym. Sci., 92, 543 (2004).124. S. Varghese and J. Karger-Kocsis, Polymer, 44, 4921 (2003).125. J. H. Lee, D. Jung, C. E. Hong, K. Y. Rhee and S. G. Advani, Compos. Sci. Technol., 65, 1996 (2005).126. Y. C. Ke and P. Stroeve, Polymer-Layered Silicate and Silica Nanocomposites, Elsevier (2005).127. W. Liu, S. V. Hoa and M. Pugh, Composites Science and Technology, 65, 307–316 (2005).128. Y. X. Zhou, P. X. Wu, Z-Y. Cheng, J. Ingram and S. Jeelani, eXPRESS Polymer Letters 2, 1, 40–48 (2008).129. A. Lee and J. D. Lichtenhan, Journal of Applied Polymer Science, 73, 1993–2001 (1999).130. J. T. Kim, T. S. Oh, and D. H. Lee, Polym. Int., 52, 1058 (2003).131. T. P. Mohan, M. R. Kumar and R. Velmurugan, J Mater Sci., 41, 2929–2937 (2006).132. L. F. Cai, Y. L. Mai, M. Z. Rong, W. H. Ruan and M. Q. Zhang, eXPRESS Polymer Letters, 1, 2–7 (2007).133. J. C. Halpin, J. Compos. Mater. 3, 732 (1969).134. J. C. Halpin and J. L. Kardos, Polym. Eng. Sci., 16, 344 (1976).135. T. Mori, K. Tanaka, Acta Metall. Mater., 21, 571 (1973).136. Y. Benveniste, Mech. Mater., 6, 147 (1987).137. T. Mura, Micromechanics of Defects in Solids, 2nd ed., Martinus Nijhoff, Boston (1987).138. G. P. Tandon and G. J. Weng, Polym. Compos., 5, 327 (1984).139. D.A. Brune and J. Bicerano, Polymer, 43, 369 (2002).140. T. D. Fornes, D. R. Paul, Polymer, 44, 4993 (2003).

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