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Introduction

CHAPTERl

Introduction

Chapter I

Introduction

1.1 Polymers for space applications

Chapter I

Polymers used in space application should possess good thermal stability,

improved toughness, strength, out gassing and barrier property because of the

harsh environments of the space such as intense radiation, extreme temperature

excurSIOns, vacuum, atomic oxygen and high velocity impact from

micrometeorites. High performance polymers like epoxy, polyimide, phenolic

resin, polyether ether ketone, polyether sulfone, cyanate ester etc have great

potential in space applications owing to their excellent thermal and mechanical

properties and chemical stability. [loAJ

1.1.1 Phenolic resin

Phenolic resins are large family of polymers and oligomers, composed of a

wide variety of structures based on the reaction products of phenol or substituted

phenol with an aldehyde, especially formaldehyde, in the presence of an acidic or

basic catalyst. Phenolic resins are widely known as phenol -formaldehyde resins,

PF resins and phenoplasts. Reaction of phenol with formaldehyde involves a

condensation reaction, which leads to a cross-linked polymer structure. [1-4].

Phenolic resins are employed in a wide range of applications, from

commodity construction materials to high technology applications in electronics

and aerospace. Even though phenolic resin is not an exact substitute for epoxies

and polyimides in many engineering applications, their composites find a major

demand in thermo structural application in aerospace industry including aircraft

interiors and rocket nozzles. This is due to good heat and flame resistance,

excellent ablative properties and low cost of the phenolic resin [5-7].

1.1.2 Cyanate ester

Cyanate ester belongs to a unique class of thermosetting resins falling in

high performance criteria and is currently gaining industrial applications. Cyanate

ester resin is produced from the reaction between a cyanogen halide and a phenol

in the presence of a base. Cyanate ester shows high glass transition temperature in

the range of 290-300°C. They possess good toughness, low dielectric constant and

2

Introduction Chapter 1

dissipation factor, radar transparency and low moisture absorption, which make

them a resin of high performance structural and thermo mechanical applications in

aerospace industry and electronic applications [8-11]. The low out gassing,

minimal dimensional changes during thermal cycling, good long term stability, self

adherent properties to honeycomb and foam cover, good electrical properties and

high service temperature are the key advantages of cyanate resin over the state of

the art epoxy resin [12, 13j 13~

1.1.3 Polyether ether ketone (PEEK)

Polyether ether ketone (PEEK) is a thermally stable, high performance,

semicrystalline thermoplastic having outstanding mechanical, thermal and

chemical properties like toughness, stiffness, thermo oxidative stability, chemical and

solvent resistance, electrical performance, flame retardancy and retention of physical

properties at high temperature [14].

Rigid rod structure of the PEEK polymers is disadvantageous for many

applications. One of the common methods to reduce the rigidity of the polymer is

by introducing meta linkages or flexible spacers such as ether, thio, propylidene

linkages into the polymer chain. The introduction of bulky pendant groups into the

polymer backbone [15- 17] will improve the solubility and thermal properties of the

polymer. Rao and Sivadasan [18] have synthesised a series of PEEK polymers by

varying the reaction solvents, temperature and monomers. PEEK was synthesized

by nucleophilic substitution reaction of 4,4'-difluorobenzophenone/4,4'

dichlorobenzophenone with hydroquinone, resorcinol and dihydroxybenzophenone

in various solvents using anhydrous potassium carbonate as catalyst. All the

polymers showed good thermal stability. In another work, alkyl substituted PEEK

were synthesized to reduce the crystallinity and melting temperature to increase the

glass transition temperature to facilitate the ease of processability at relatively low

temperature [19].

Poly ether ether ketone (PEEK) [20, 2 I] is insoluble in all common organic

solvents at room temperature except for self protonating acids such as H2S04, HF,

3

Introduction Chapter I

methane sulfonic acid and trifluromethane sulfonic acid. In addition, solubility of

PEEK was found in benzophenone above 220°C and in chloronapthalene above

230°C [22-24]. Due to its balanced properties including good strength, high

ductility, high and low temperature stability (+250 to -250°C), excellent impact

strength and excellent resistance to moisture, liquid oxygen (LOX) and solvents,

PEEK is being used in a wide range of applications such as electrical and

electronic parts, military equipments, automobile parts, wires and cables in nuclear

plants, underground railways and oil wells, as well as advanced structural

composites for aircrafts and as a seat material for cryo valves in cryo engines.

Carbon / PEEK is an ideal candidate for Liquid Hydrogen (LH2) and LOX cryo

tanks and it has been used in HOTOL re-entry vehicle. They are also being used in

the helicopter, cryogenic tank lining, pressure vessels, rocket motor, solar panel

substrate, antenna reflectors and sail boom L24-tiI·

Francis et al [25-31] have synthesized a senes of oligomers having

molecular weight of M n = 3000 to 12000 g/mol such as hydroxyl terminated poly

ether ether ketone having pendent ditert-butyl, tert-butyl and methyl groups and

poly(ether sulfone ether ketone) in order to make use of them as a toughening

agents for epoxy resin. They were synthesised by the nucleophilic substitution

reaction of 4,4' -difluorobenzophenone with 2,5-ditert-butylhydroquinone, tert

butylhydroquinone, methylhydroquinone and 4,4' -dihydroxydiphenylsulfone

respectively using potassium carbonate as catalyst in N-methyl-2-pyrrolidone

/sulfolane medium. The molecular weight was controlled by taking required

amount of monomers using Carother's equation [32,33]. The introduction of the

pendant alkyl groups reduced the crystallinity of the PEEK polymer. The ditert

butyl PEEK was found to be more crystalline compared to the other polymers due

to the symmetrical structure of the former. The Tg of the modified PEEK polymers

was dependent on the molecular weight and polymer structure. Ditert-butyl PEEK

polymers showed melting peak due to semicrystalline behaviour. PESEK showed a

Tg of 192°C due to the presence of polar sulfone groups in the polymer chain.

Among the tert-butyl and methyl PEEK polymers, tert-butyl PEEK polymers

4


showed higher Tg due to the presence of bulky pendent tert-butyl group. The

polymers were found to be thermally stable up to 450°C in nitrogen atmosphere.

All the polymers except ditert-butyl PEEK polymers were miscible in epoxy resin

before curing.

1.1.4 Polyether sulfones

Polyether sulfone is an important class of engineering thermoplastic. They

are amorphous polymers with excellent thermal and mechanical properties. Inspite

of their excellent properties such as high Tg and superior thermo-oxidative stability

[34], some undesirable features like solvent sensitivity and creep under load at

elevated temperatures (> 175°C) [34,35] preclude their use when solvent resistance

and high temperature dimensional stability are required. However compared to

poly aryl ether ketones they are considerably more fluid at high temperatures and

thus are moulded and processed more easily [36]. Poly sulfone (bisphenol-A type)

was commercialized by Union Carbide \'lith the trade name Udel in 1966.

Polyether sulfone (bisphenol-S type)] "vas commercialized by ICI with trade name

Victrex PES in 1971; poly phenyl sulfone (bisphenol type) was produced

commercially by Union Carbide with trade name Radel in 1972. The first synthetic

procedure was based on the condensation of alkali metal salts of various

bisphenols with activated halogeno aromatics such as 4, 4' -dichlorodiphenyl

sulfone (OCOPS) [37]. In nucleophilic substitution reactions, the rates are

dependent on the basicity of the bisphenol salts [38] and on electron withdrawing

power of the activating group in the dihalide [37,39].

Polyether sulfone (PES) is also inherently flame retardant for use in

electronic components and testing devices. The incorporation of sulfone groups

into the macromolecular chain often generates interesting properties such as

increased Tg and higher decomposition temperature and can also lead to liquid

crystallinity due to strong interactions between sulfone groups [40]. These

polymers retain their useful mechanical properties between the temperature limits

of-100°C to 175°C [41]. PES is resistant to attack by non-oxidizing acids, alkalies,

salts and aliphatic hydrocarbon solvents such as gasoline. They dissolve in

5


chlorinated hydrocarbons (CCI4) and are attacked by concentrated H2S04. They

swell in aromatic hydrocarbon solvents, benzenes, esters and ketones. The solvent

resistance of PES can be increased by introducing crystalline polymer blocks such

as PEEK, PEK, or by the introduction of cross-links in linear polymer systems by

appropriate chain-end functionalization [42-46]. Owing to their excellent thermal

index of 160°C, creep resistance, transparency and mechanical and electrical

properties, poly sulfones are used as printed circuit boards, integrated circuit

carriers, television and stereo components, aircraft components, refrigerator parts,

filtration membranes, pipes etc. Poly ether sulfone can also be used as a

toughening agent for epoxy resin.

1.1.5 Polyimide

Polyimides have the characteristic functional group of

/CO.JVV\\lVV'N

'CO.JVV\, with a back bone that consists predominantly the ring

structures and hence high glass transition temperature. Furthennore, many of the

structures exhibit a high level of thermal stability and can be used for high

temperature applications.

The general method of preparation of polyimide is shown in Fig 1.1. In the first

step polyamic acid is formed by the reaction of aromatic diamine with an aromatic

tetra carboxylic acid dianhydride in the presence of dipolar aprotic solvent at RT.

In the second step, polyamic acid will be converted into polyimide either by heat

treatment or chern ical treatment [47].

Pyromellitic dianhydride (PMDA) is the most commonly used dianhydride. A

number of diamines such as m-phenylene diamine, benzidine and oxy dianiline

(ODA) have been found to give polyimide with high degree of oxidative and

thermal stability. The commercially available Kapton (DuPont) polyimide film is

derived from PMDA and ODA. The usual solvents employed for the preparation of

polyamic acid include dimethyl formam ide and dimethyl acetamide [1.47].

Aromatic polyimides possess excellent mechanical strength and thermal stability.

During the past decades, research interest in these polymers has been increased in

6


view of the increasing demand for technological applications in a variety of fields

such as aerospace, automobile, and microelectronics [1, 47,47~J

n

o 0

II IIHO-C C-NH-Ar'

""A//r"

HN-C C-OH

II IIo 0

Polyamic acid

RT+ H2N--At-----NH2

dipolarAromatic diamine aprotic

solvent

o 0

II II/c"" /c""

O""C/Ar"c/O

II IIo 0Aromatic tera carboxylicdianhydride

~Heat

nPolyimide

Fig1.! General scheme for the preparation of polyimide

1.1.6 Epoxy resin

The term 'epoxy resin' refers to both the prepolymer and its cured resin/

hardener system. The former is a low molecular weight oligomer that contains one

or more epoxy groups per molecule (more than one unit per molecule is required if

the resultant material is to be crosslinked). The characteristic group, a three

membered ring known as the epoxy, epoxide, oxirane, glycidyl or ethoxyline group

is highly strained and therefore very reactive. Epoxy resins can be cross-linked

through a polymerization reaction with a hardener at room temperature or at

elevated temperatures (latent reaction). Curing agents used for room temperature

cure are usually aliphatic amines, and for high temperature, aromatic amines and

acid anhydrides are used. Polyfunctional amines, polybasic carboxylic acids,

mercaptans and inorganic hardeners are also used as specialized curing agents. In

general, the high temperature cured resin systems have improved properties, such

as higher glass transition temperature, strength and stiffness, compared to those

cured at room temperature [49-53].

7


Among the thermoset materials, epoxy resins shows special chemical

characteristics such as absence of byproducts or volatiles during curing reactions,

low shrinkage up on curing, curing over a wide temperature range and the control

of degree of cross-linking. Depending on the chemical structure of the curing

agents and curing conditions, the properties of cured epoxy resins will vary. Epoxy

resins are versatile with excellent chemical and heat resistance, high adhesive

strength, good impact resistance, high strength and hardness, and high electrical

insulation [51-53].

Fig 1.2 illustrates the cure reaction of an epoxy resm with an amme

hardener [48]. The two different functional groups react during the initial

conversion (Reaction I) and form a linear or branched polymer. The addition of the

primary amine to an epoxide group leads to the formation of a hydroxyl group and

a secondary amine, which continues until the primary amine groups are exhausted.

Reaction II illustrates the cross linking through the addition of secondary amines

with epoxy groups, where the macromolecules develop a three-dimensional

network. One of the most common side reactions is etherification (Reaction III),

where a hydroxyl group reacts with an epoxide group, forming an ether linkage

and a further hydroxyl group. The extent to which etherification takes place during

cure depends on the structure and chemistry of the resin and hardener, as well as

the cure conditions. When the branched structures extend throughout the system,

the gel point is reached. At this characteristic point, the cross-linked resin does not

dissolve in any solvent, although a soluble (sol) fraction may still be extractable.

Further, diffusion-controlled cure is required to increase the degree of cross-linking

to produce a structural material with a high modulus of a vitrified or glassy solid

material. After the completion of cure, epoxy resin attains highly cross-linked

structures, which have good mechanical and thermal properties. In addition they

have very good adhesive properties. The properties of the cured resins are

dependent on the curing agent. Amine cured epoxy resins have excellent electrical

properties coupled with excellent chemical resistance and relatively high heat

distortion temperature (HOT). Epoxy resins cured with anhydrides have excellent

mechanical properties with better outdoor weathering resistance. Some of the

8


anhydride-cured resins have high HDT and excellent retention of strength at high

temperature. These systems have low viscosity and long pot life [51, 53].

+ H~\-/CH-R' __..~

o

CH2-CH-R'

R-N/ I (I)

I OH

H

CH2-CH-R' +

R-N/ II OH

H

H2C\-/-CH-R" ---l"~

o

/CH2-CH-R'

R-N I (II)

'" OHCH -CH-R"2 I

OH

R-CH-R'

IOH

Fig 1.2 The three possible reactions during cure of an epoxy resin with an amine(I) primary amine-epoxy addition, (II) secondary amine-epoxy addition, (III)etherification

1.2 Toughening of epoxy

Epoxy resins are extensively used in wide variety of applications in the

field of aerospace, automotive, electronic and constructive industries because of

their outstanding mechanical and thermal properties, ease of processing and

chemical resistance. All these properties are attributed to their high cross-link

density. Unfortunately owing to the high cross-link density, epoxy resins have very

low resistance to crack initiation and propagation (poor fracture toughness). It is

very important to increase the fracture toughness of these materials without

causing significant decrease in the mechanical properties. General strategy widely

used to improve the fracture toughness of the epoxy resin was the addition of

second phase, which will separate during the curing of epoxy and produce a multi

phase morphology, which enables to initiate a variety of toughening mechanisms

during the crack propagation. Rubbers, Inorganic glasses, thermoplastics etc have

9


been employed to increase the fracture toughness of epoxy resin. One of the

effective methods to increase the fracture toughness of epoxy resin is by the

addition of elastomer (rubber), but it often resulted in the reduction of modulus and

high temperature mechanical properties [54-55]. On the other hand, toughening of

epoxy with high performance engineering thermoplastic improves the fracture

toughness with the retention of modulus and high temperature mechanical

properties. This is due to tough, ductile, chemically and thermally stable and high

Tg characteristics of the thermoplastic.

1.2.1 Thermoplastic toughening of epoxy resin

Extensive study on the toughening of epoxy with engineering thermoplastic

such as PEEK and PES has been reported in this section.

1.2.1.1 Poly (ether ether ketone) toughened epoxy resin

According to Zhong et al [56] epoxy/phenolphthalein poly(ether ether ketone)

(PEK-C) blends cured at 80aC using DDM as the curing agent are found to be

homogeneous as evident from DSC, DMA and SEM studies. In contrast to this

when the temperature is raised to 150a C, the blends became heterogeneous [57].

Scanning electron microscopy and DMA studies confirmed the phase separated

morphology. The difference in phase behaviour is due to the increased mobility of

the molecules at high curing temperature. In order to establish the relationship

between the curing agent and phase behaviour, Guo et al [58] investigated blends

cured with maleic anhydride (MA),Phtaleic anyhydride (PA) and

hexahydrophthalic anhydride (HHPA). Phase separated morphology is evident

from SEM micrographs for MA and HHPA cured systems. But due to the aromatic

nature of PA and PEK-C, the cured blends are homogeneous.The mechanical

properties of DDM cured PEK-C/epoxy blends revealed that fracture toughness

and flexural properties are slightly lowered by the addition of PEK-C. The

reduction in properties was due to the homogeneous nature of the blends.

Bennet et al [59] have prepared the epoxy/poly aryl ether ketone blends

with a series of amine terminated oligomers based on tert-butyl hydroquinone

10


(TBHQ) and methyl hydroquinone (MeHQ). All the blends showed improved

toughness. They observed 350-750% increase in the fracture energy for the blends

with either no loss or a slight increase in modulus compared to neat epoxy/amine

resin. The substantial increase in fracture energy with TBHQ based oligomer was

strongly dependent on the final morphology of the blends. The phase inverted

morphology gave maximum toughness. The modification of epoxy with amine

terminated PEEK was reported by Cecere and Mc Grath [60]. The water absorption

reduced by half by the addition of 40% oligomer. lijima et al [61] effectively used a

series of synthesized poly (aryl ether ketone)s to improve the fracture toughness of

DGEBA resin with no expense of mechanical properties. They obtained maximum

toughness by the addition of 10wt% modifier. Cocontinuous morphology gave the

maximum toughness. Song et al [62] used PEK-C to modify a tetrafunctional

epoxy (TGDDM). Brostow et al [63] used fluorinated poly aryl ketone (12F-PEK)

to reduce the friction of DGEBA resin cured with triethylenetetramine (TETA).

Addition of 10% fluoropolymer, reduced friction by 30%. They also found that the

properties were strongly dependent on the curing temperature. They further observed

that 5% addition of 12F-PEK to epoxy cured at 24°C resulted in better healing after

scratching. The properties were dependent on the cure temperature and hence can be

manipulated by changing cure temperature [64]. The total surface tension of a polymer

solid has strong effect on the static friction, dynamic friction, penetration depth and

residual scratch [65].

Francis et al [25-31] has studied the toughening of epoxy using hydroxyl

terminated PEEK having pendent ditert-butyl (PEEKDTOH), tert-butyl

(PEEKTOH) and methyl (PEEKMOH) groups and poly (ether sulfone ether

ketone) (PESEKOH). The advantage of functionally terminated PEEK with pendent

alkyl groups instead of virgin PEEK to toughen the epoxy resin is that the oligomers

are less crystalline and easy to process. All blends show good miscibility with epoxy

resin except ditert-butyl modifed PEEK. Triphenylphosphine (TPP) is used to

catalyse the reaction between hydroxyl groups in the PEEK polymer and epoxy

functional group of DGEBA epoxy resin. All the blends are cured by using DDS

and the cure kinetics revealed an autocatalytic reaction. The blends become opaque

II


on curing with 0 OS as a result of reaction induced phase separation (RIPS). Well

defined spherical domains of PEEK distributes uniformly in the epoxy matrix are

obtained for tert-butyl and methyl PEEK toughened epoxy resin, whereas non

spherical domains are obtained in ditert-butyl PEEK modified epoxy resin. In

15phr PEEKM modified epoxy resin, a complex morphology is observed. The

change in morphology is due to high viscosity of epoxy/15phr PEEKM blend as a

result of its high molecular weight. The rate of phase separation decrease as a

consequence of high viscosity compared to 5 and 10phr PEEKM modified epoxy

resin. Phase separation in the blends is expected to occur by nucleation and growth

mechanism. The domain size and its distribution are dependent on the composition

and molecular weight of the polymer. Hydroxyl terminated polymers give smaller

domains compared to random polymers due to the chemical interaction between

the hydroxyl terminated polymers and epoxy resin. The cure kinetics of OGEBA

epoxy modified with PEEK based on tert-butyl hydroquinone (PEEKT) is found to

follow autocatalytic mechanism. The rate of reaction decreases with the addition of

PEEKT to epoxy resin [27]. PESEK modified epoxy resins reveal no heterogeneity

in the blends. The homogeneous morphology of the blends is due to structural

similarity of PESEK with cured epoxy resin and H-bonding interaction between

PESEK and epoxy resin. All the thermoplastics are effective in improving the

fracture toughness of epoxy resin. The extent of improvement is dependent on the

type and molecular weight of the modifier and composition of the blends. Tert

butyl and methyl PEEK polymers are the most effective modifiers among the

different systems studied. In PESEK toughened epoxy resin, the increase in

fracture toughness is less compared to tert-butyl and methyl PEEK toughened

epoxy resin due to homogeneous morphology.

1.2.1.2 Polyether sulfone modified epoxy resin

Blends of epoxy resin with polyether sulfones were the most studied among

the various thermoplastic modified epoxy resins. Toughening of epoxy with

polyether sulfone (PES) was first studied by Bucknall and Partridge [66]. They

observed no improvement in fracture toughness with the addition of PES to a tetra

12


functional epoxy resin. Jenninger et al [67] studied the cure kinetics of DDM cured

PES toughened DGEBA epoxy resin. Cure kinetics of the blends was found to

follow autocatalytic reaction. Decrease in cure temperature or increase in

thermoplastic lead to less perfect cross-linked structure and hence low Tg• The

influence of trifunctional epoxy on the cure kinetics of DGEBAIPES blends cured

with MCDEA was investigated in detail by Alig et al [68]. PES with terminal reactive

phenoxy group was found to accelerate the reaction rate.

Kim and Inoue [69] used dynamic mechanical analysis (DMA) and FTIR to

follow the cure reaction during the late stages of DGEBAIPES/DDS system. They

observed the increase in Tg with cure time. Two Tg's were obtained from the

dynamic mechanical spectrum of the blends. Yamanah and Inoue [70] studied the

phase separation mechanism in PES toughened amine cured epoxy resin. From

light scattering studies, these authors found that phase separation occurred by

spinodal decomposition. Swier and Mele [71, 72] used modulated temperature

differential scanning calorimetry to study the phase separation mechanism in PES

modified epoxy resin. They reported that careful selection of cure schedule will

lead to PES modified epoxy with tailor made morphologies. They were able to

detect phase separation earlier than the conventional cloud point measurements. In

DDM crosslinked system, complex morphology with interphase at certain cure

schedules was obtained. The heat capacity signal for 20 phr PES/DGEBA/DDM

system showed two step decrease; first one was due to phase separation of PES and

the second one due to vitrification of epoxy [73]. Guo [74] reported that the ultimate

morphology of phenolphthalein poly (ether sulfone) (PES-C)/epoxy blends was

dependent on the type of curing agent used. Amine cured (DDS and DDM) blends

gave homogeneous morphology because of H-bonding whereas phthalic anhydride

(PA) cured blends gave heterogeneous systems. The reaction mechanism was step

growth polymerisation in amine curing and chain growth in anhydride cured blends. In

poly(ether sulfone):poly(ether ether sulfone) (PES:PEES) copolymer/epoxy resin

system, MDEA cured blend was heterogeneous while 3,3'-DDS cured blend was

homogeneous.The difference in morphology was due to different viscosity and

solubility effects of the blends [75,76]. But according to Bucknall and coworkers

13


[77], molar mass of epoxy resm was the mam factor in controlling the phase

separation in epoxy/PES blends. Alig et al [78] studied the RIPS in PES modified

epoxy resin using SALS. Ratna et al [79] synthesised amine terminated polyamide

sulfone (ATPS) and blended it with epoxy resin. The HY 951 cured resin formed

homogeneous blends before and after curing.

The morphology of OGEBA/PES blends was also dependent on the harder

to epoxy ratio [80]. In OOM cured blends with lower stoichiometric ratio,

spherical domains were observed and at stoichiometric ratio co-continuous

morphology was observed. Above stoichiometric ratio, interconnected globular

structure was observed. Pasquale et al [81] used amine functionalised PES to

increase the fracture toughness of epoxy resin. The improvement in fracture

toughness was due to the decrease in crosslink density. PES with pendent amino

group and vinyl benzyl groups were used to modify epoxy resin [82,83]. The

amount of amine groups in the polymer was critical in enhancing the fracture

toughness and morphology. The presence of dicumyl peroxide (OCP) gave more

fracture toughness compared to benzyl-PSF modified epoxy resin. Iijima et al [84]

reported that chloro terminated PES was more effective in enhancing the fracture

toughness than epoxy terminated PES. The increase in toughness was due to plastic

deformation of matrix prior to failure. Huang et al [85] reported homogeneous

blends of epoxy and bisphenol-A PSF cured with OOM having 20% increase in

fracture toughness. The morphology and properties of bisphenol-A PES modified

biphenyl type epoxy resin were dependent on the curing conditions employed [86].

Lower curing temperature favoured homogeneous morphology and higher cure

temperature favoured heterogeneous morphology. In both cases fracture toughness

increased without decreasing other mechanical properties. In phase separated

systems the increase in fracture toughness was due to crack branching and in

homogeneous blends the increa.se was due to induced shear deformation of the

matrix. The addition of PES and coreshell rubber particle with polybutadiene core

and poly(methyl methacrylate) (PMMA) shell increased the fracture toughness of

OGEBA epoxy resin cured with DOS. The toughening mechanism was rubber

cavitation followed by plastic deformation of the matrix [87].

14


Mimura and Ito [88] found that rate of decomposition in PES modified

biphenyl epoxy cured with phenol novolac resin was dependent on the morphology

of the blends. Co- continuous blends decomposed fast. The physical ageing in

PES/epoxy/DDS blends was also investigated [89, 90]. Raman spectroscopy was a

useful tool for studying the reaction between epoxy resin and amine terminated

copolyethersulphone. Reaction between the blend components was confirmed in

the absence of curing agent [91]. The curing reaction of TGAP/0OS system

decreased with increase in PES content [92]. From calorimetric studies it was

found that the curing exotherm decreased linearly with increase in PES content.

The rheological properties during curing were investigated and found that PES

delayed the curing reaction [93]. Hydroxyl, amine and non-functional PES /TGAP

blends cured with DDS showed two phase morphology [94, 95]. The two-phase

morphology in MDEA cured PES modified epoxy resin was investigated using

micro-Raman spectroscopy [96]. Gumen and Jones [97] used group interaction

model to examine the contribution of individual components to thermo-mechanical

properties of the blend.

A mixture of TGAP and diglycidyl ether of phenol-formaldehyde in the

ratio 2:1 was modified with hydroxyl end capped PES and studied the morphology

by curing with 3,3'-DDS. Co-continuous morphology was observed at 30 and 40%

modifier and phase inverted morphology was observed at 50% PES. Co-continuous

morphology facilitated more uniform stress distribution under load and avoids

premature failure [98]. Raghava [99,100] obtained bimodal particle size

distribution for MY nO/hydroxyl PES blends. The properties of the blends are also

dependent on the type of PES used to modify epoxy resin. On changing from non

functional PES to functional PES, the particle size decreased. On adding 20 to 30

phr functional PES, the morphology changed from co-continuous to phase inverted

one and maximum toughness was observed in this region [101]. Incorporation of

triepoxy to diepoxy was found to be more effective than changing the molecular

weight of PES. A full interpenetrating polymer network (lPN) with increased

fracture toughness was obtained on adding TGDDM to diglycidyl ether of

bisphenol-S (DGEBS)/PES:PEES copolymer blend cured with DDS. Tg was also

15


increased with the addition ofTGDDM [102]. Francis et al [31,103] has studied on

the PES toughened epoxy/DDS blend. The blends were homogenous in nature due

to the H-bonding and also the blends might have gelled before phase separation

due to high blend viscosity and fast curing reaction. The improvement in fracture

toughness was found to be less than the PEEKMOH or PEEKTOH toughened

system [25-31] due to the absence of phase separation.

1.3 Toughening mechanisms in thermoplastic modified epoxyresins

The optimum properties for thermoplastic modified epoxy reSIn were

strongly dependent on the intrinsic properties of the component materials. In

general several toughening mechanisms operate simultaneously to produce the

overall toughening effect. The important mechanisms responsible for enhancement

of fracture toughness are given below.

1.3.1 Thermoplastic-particle bridging

Thermoplastic particles span the crack, which necessitates their ductile

stretching and tearing. This mechanism provides closure traction to the crack

surfaces and effectively reduces the local stress intensity factor at the crack tip. It

may induce significant toughening effect especially for highly cross-linked

epoxies, owing to the high yield stress of the thermoplastic modifier. This

mechanism is effective only when relatively large particles with strong interfaces

are present.

1.3.2 Crack pinning or bowing by thermoplastic particles

Rigid thermoplastic particles act as impenetrable objects and effectively pin

the advancing crack. The pinned crack front bows out, which consumes additional

energy. This mechanism is effective only when there is good interfacial adhesion.

1.3.3 Crack deflection and/or bifurcation by thermoplastic particles

Thermoplastic particles will change the crack path by deviating the crack

from its principal plane of propagation and/or by splitting into several secondary

16


cracks. The deflection of the crack path increases the total surface area of the crack

surface. The bifurcation of a crack reduces the local stress intensity factor at the tip

by distributing it over multiple cracks. This mechanism accompanies other

toughening mechanisms.

1.3.4 Shear banding of the matrix

This mechanism provides significant increase in fracture toughness if the

intrinsic ductility of the matrix is fully exploited. The thermoplastic acts as stress

concentrators owing to the significant modulus mismatch between the particle and

the matrix. This stress concentration causes the matrix to undergo extensive shear

deformation, generally in the form of massive shear banding between the particles.

The formation of shear bands absorbs considerable energy, thereby increasing the

fracture toughness.

1.3.5 Mirocracking of the matrix

Thermoplastic particles causes stress concentration and initiate massive

microcracks in the surrounding matrix, dissipating extra energy and increasing the

fracture toughness. This mechanism is effective for highly crosslinked epoxies

when the particles are relatively rigid and capable of dehonding.

1.3.6 Transformation of thermoplastic particles

Particles of semicrystalline thermoplastics may undergo stress induced

phase transformation in crystal structure. The dilatation usually associated with this

transformation reduces the tensile stresses ahead of the crack tip, thus effectively

lowering the local stress intensity factor. This mechanism is valid only for epoxies

toughened with semicrystalline thermoplastic.

1.4 Morphology of the thermoplastic toughened epoxy blends

Most of the thermoplastics are miscible with epoxy resin in the uncured

state, but on curing the miscible binary epoxy/thermoplastic blends can remain as

miscible blend or phase separated system depending on the extent of interactions

17


between the blend components, cunng agent and cunng conditions. But

thermoplastics like polyamide 12 (PA12) and poly(vinyledene fluoride) [104,105]

are immiscible with epoxy resin and upon curing the immiscible blend invariably

resulted in heterogeneous blends. In the absence of strong interactions, phase

separation occurred in the blends. The different morphologies developed in epoxy

resin/thermoplastic blends are shown in Fig. 1.3.

\EpoXy' resin + Thermoplastic I

Mixing

I I

• .-..• • •• • • ••••••• • ••

I..- ..• • •• ••••

•••••• ••Heterogeneous

Heterogeneous

\ Curing

Heterogeneous

Homogeneous

Homogeneous

\ Curing

I I

• .-.•• • •• • • ••

••••• • ••

Fig 1.3 Different types of morphologies in epoxy resin/thermoplastic blends

From the above discussions on thermoplastic modified epoxy resin, it was

found that the ultimate properties of the blends were strongly dependent on the

final morphology. Most of the blends undergo reaction induced phase separation

(RIPS) upon curing. The main advantage of RIPS is that different morphologies

can be obtained by changing the composition, curing agent and curing conditions.

A uniform dispersion of rubber or thermoplastic can be obtained by RIPS. Even

homogeneous morphology was obtained by proper control of curing conditions.

Depending on the composition, various morphologies are generated as shown in

Fig 1.4.

18

Introduction

Homogeneous blend

Phase s\parationduring curing

/1'

Chapter I

-. .. . .•• ••• • •••• • • • ••• ••• ••

dispersed thermoplastic co- conti nu ou s dispersed epoxy(phase-inverted)

Fig 1.4 Schematic diagram showing the evolution of morphology during curing forepoxy resin/thermoplastic blends

1.5 Phase separation mechanism in blends

Most of the miscible. polymer blends form single phase system over a

certain temperature range. The occurrence of phase separation at a given

temperature is determined by the shape of the free energy of mixing (~Gm) versus

composition curve. Fig 1.5 shows a typical LCST curve. The composition of the

homogeneous mixture is given by Co, C'aand C"a are the concentrations ofthe nuclei.

The solid line represents the binodal curve and the dashed line represents the spinodal

curve. Phase separation can occur either by spinodal decomposition or by nucleation

and growth mechanism depending on the blend composition and temperature [106].

19

Introduction

,.,E'""""e

LL.

o

\\\, I

a c~ c~

Volume fraction of component 2

(b)

(a)

Chapter I

Figure 1.5 (a) LeST curve and (b) free energy of mixing versus composition for abinary polymer mixture [106]

In the spinodal region, a small fluctuation in composition will lead to phase

separation due to decrease in free energy. There is no thermodynamic barrier to phase

growth and the composition of the phases changes continuously with time. Spinodal

decomposition leads to co-continuous structure. The evolution of phase growth and

corresponding phase structure is schematically represented in Fig 1.6.

In the region between spinodal and binodal points, the blend is in the meta

stable state, because small concentration fluctuation does not lead to phase

separation. At high concentration fluctuation, demixing occurs by nucleation and

growth mechanism and the corresponding phase structure is shown in Fig 1.7.

20

Introduction

c~ ------------------

(a)

Chapter I

(b)

Figure 1.6 Phase separation in miscible polymer blend by spinodal decompositionmechanism (a) one dimensional evolution of concentration profiles (b) resultantphase structure [106]

c:

Co

c~

(a)

(b)

Figure 1.7 Phase separation in a miscible polymer blend by nucleation and growthmechanism (a) one dimensional evolution of concentration profiles (b) resultantphase structure [106]

21

Introduction

1.6 Polymer clay nanocomposites

Chapter I

Polymer nanocomposites are generally defined as the combination of

polymer matrix and fillers that have at least one dimension in nanometer range.

The nano fillers can be one- dimensional (layered minerals such as clay) [107

114], two-dimensional (like carbon nanotube, nanowires, nanofibers, cellulose

whiskers etc) [115-117], and three-dimensional (spherical particles include silica

nanoparticles, nanowhiskers etc) [118-121]. Polymer nanocomposites are known

for its outstanding mechanical properties like high elastic modulus [107-112,122

125], increased strength [107-112,126], barrier resistance [127-131], flame

retardancy [132-139] etc with very small addition (:S 5 wt%) of nano particles. This

is due to the very large surface area of interaction between polymer matrix and

nano filler. Among the different nanofillers, special attention has been paid to clays

in the field ofnanocomposites. Layered silicates (clays) are found to be one of the

ideal nano reinforcements for polymers, because of its high intercalation

chemistry, high aspect ratio, ease of availability and low cost.

1.6.1 Clays

Clays are hydrous silicates or alumino silicates and are fundamentally

containing silicon, aluminum or magnesium, oxygen and hydroxyl with various

associated cations. These ions and OH groups are organized into two dimensional

structures as sheets. Clay minerals are also called layered silicates or phyllo

silicates because their structural frame work is basically composed of 1 nm thick

silicate layers comprising silica and alumina sheets joined together in various

proportions and stacked on top of each other in certain way with a variable

interlayer distance. The clay minerals can be classified into three different types

based on the condensation ratio of silica to alumina sheet [113,114].

1:1 type: It consists of one octahydral alumina sheet and one tetrahydral silica

sheet condensed in I: 1 ratio, called as dimorphic or two sheet minerals. The sheet

does not bear any charge due to the absence of isomorphic substitution in both

silica and alumina sheet. Layers are held together by hydrogen bonding between

22


hydroxyl group in octahedral sheets and oxygen in tetrahedral sheets. The space

between the layers is occupied by water molecules. Eg: Kaolinite, Perlite,

Hallosite etc:

2:1 type: In 2: I, trimorphic or three sheet minerals, an alumina sheet is

sandwiched between two silica sheets. This type of clay belongs to smectite

family. Stacking of these layers create a Vander Waals gap between clay layers.

Isomorphic substitution of A13+ with Fe2+, Mg 2+, Li+ in the octahedron sheets

and/or Si4+ with A1 3+ in the tetrahedron sheets gives each layer an overall negative

charge, which is counterbalanced by exchangeable metal cations such as Na+, Ca2+,

Mg2+, Fe2+ and Lt residing in the interlayer space. Eg: montmorillonite (MMT),

hectorite, saponite, tlouro hectorite,laponite,tlouromica (somasif)

2:2 type: The layer structure of 2:2 type, known as tetramorphic of four sheet

minerals, is formed by the alternate condensation of silica tetrahedron sheets and

alumina or magnesium octahedron sheets. Eg: Chlorite

1.6.1.1 Montmorillonite

Among the different types of clay minerals, montmorillonite is the most

commonly used for the preparation of polymer clay nanocomposites [107-114].

Montmorillonite owes special attention among the smectite group due to its ability

to show extensive inter layer expansion or swelling, because of its peculiar

structure as shown in Fig 1.8.

The crystal structure of montmorillonite consists of layers formed by

sandwiching an edge shared octahedral sheet of alumina between two silica

tetrahedral sheets, so that the apical oxygen atoms of the tetrahedral sheets are all

shared with the octahedral sheet. Isomorphous substitution of aluminium for

silicon in the tetrahedral sheet and iron or magnesium for aluminium in the

octahedral sheet provides an overall negative charge to the crystal lattice. As the

surface between the layers is n~gatively charged it attracts cations such as Fe2+,

Ca2+or Na l +. They form a positively charged layer between the negatively charged

surfaces of the clay layers. Such layers extend continuously in the plane constituted

23


by the x and y axes and are stacked up in the z axis, forming the whole crystal

structure. The silicate layers of MMT are planar, stiff about 1 nm in thickness with

high aspect ratio and large active surface area (700-800 m2/g). These layers

organize themselves in a parallel way to form stacks with a regular van der Waals

gap in between them, called as interIayer or gallery. The sum of the single layer

thickness and the interlayer is called d-spacing or basal spacing. The total quantity

of the absorbed cations (K+, Na+, Ca2+ and Mg2+) in the inter gallery of clay layers

at a pH value of 7 is referred as the cation exchange capacity (CEC) of clay

minerals. It is measured by the unit mill molllOOg of layered silicate (or MMT).

Higher the value of negative charge, the stronger is the capacity for hydration,

swelling and dispersion.

0 Oxygen

e Hydroxyls

• Silicon oraluminum

• Ahull.imull.iron.m8gnesil1lll

Fig 1.8 Structure of montmorillonite

1.6.1.2 Organic modification of clay

Generally clays are hydrophilic in nature. In order to make compatible with

organic polymers, the surface of the clay minerals should be modified to

organophilic prior to its use. Organic cations such as an ammonium ion or

phosphonium ion are the commonly used organic modifiers for clay minerals [107

112, 140, 14l]. Modification involves the exchange of interlayer inorganic cations

24


with organic onium salts. The organic modification causes the expansion of the

interlayer space and thereby increases the d spacing to certain extent (normally

over 2 nm). Thus the organic modification favours the diffusion of polymer or its

precursor in to the interlayer space. Fig 1.9 shows the schematic representation of

the organic modification of clay. Alkyl ammonium ions are most popular since

they can easily be exchanged with the ions situated between the layers. Depending

on the layer charge density of the clay, the alkyl ammonium ions may adopt

different structures between the clay layers. Alkyl ammonium ions reduce the

electrostatic interactions between the silicate layers thus facilitate diffusion of the

polymer into the galleries. In general, the longer the surfactant chain length, the

more will be the d spacing of the clay layers [110].

..-_.~. +H

alkylammonimn ions day organophilic clay

Fig 1.9 Organic modification of clay

1.6.1.3 Structure of polymer clay nanocomposites

Depending on the nature of the components, processing condition and

strength of the interfacial interactions between polymer and layered silicates

(modified or unmodified), either conventional composites or nanocomposites can

be formed as shown in Fig 1.1 0 [107-112]. In a conventional composite, the

polymer cannot diffuse between the clay layers and the clay particles exist in their

original aggregated state. Properties of these composites are similar to the micro

particle filled composites. An improvement in modulus is normally achieved in

conventional clay composite but this reinforcement benefit is usually accompanied

with a deficiency in other properties such as strength or toughness. If there is a

25


favorable condition for the mixing of clay minerals with polymer, two extremes of

nanocomposites are formed, i.e., intercalated and exfoliated nanocomposites. In

intercalated nanocomposites, the clay layers retain the well ordered multi structure

ofaltemating polymeric and clay layers with a d spacing of20-30 AO. On the other

hand, in exfoliated nanocomposites, the individual clay layers are well separated

and randomly distributed in the continuous polymer matrix with a d spacing of

more than 50Ao.

Phase separatedmicrocomposite

//,---'\

Exfoliated nanocomposite

Intercalatednanocomposite

Fig 1.10 Types of polymer clay nanocomposites

1.6.1.4 Preparation of nanocomposites

There are several methods to prepare clay based polymer nanocomposites.

These include in-situ polymerization, melt intercalation and solution casting.

In-Situ Polymerization: In this method, the liquid monomers are

intercalated into clay layers and polymerizes within the clay layers resulting the

expansion of the interlayer distance (d spacing). Polymerisation can be initiated by

heat or a suitable initiator. Most of the exfoliated nanocomposites are produced by

this method because it provides to select suitable reagents and polymerization

routes resulting a good affinity between clay and polymer. In situ polymersiation

technique has been used for the preparation ofnanocomposites based on polyamide

26


(PA) [131,142-146], poly methyl methacrylate (PMMA)[147-148], polY(E

caprolactone) [149], polystyrene (PS)[150], polyolefien (PP and PE) [151-154] and

polyethylene terephthalate (PET)[155]

Melt Intercalation: The melt intercalation involves the blending of clay

with the polymer matrix in molten state. If the layer surfaces have sufficient

affinity with the polymer, the polymer can diffuse between the clay layers and

form either an intercalated or an exfoliated nanocomposite [107-112,156,157].

Melt intercalation technique is used for the preparation of nanocomposites based

on polyamide [158-161], polycaprolactone (PCL) [162,163], poly lactic acid

(PLA) [164-166], poly (ethylene terephthalate (PET) [167], Polyolefins [168,169],

etc. This method is more economical and simpler than other methods. The melt

intercalation process has become increasingly popular because of its great potential

for application with rapid processing methods such as injection moulding and

extrusion. For poly methyl methacrylate/clay and poly propylene /polystyrene

clay nanocomposites, it was realized that ultrasonic assisted melt mixing

successfully generated exfoliated nanocomposites compared to In situ

polymerization [170].

Solution casting: In the solution method, polymer clay nanocomposites are

prepared by using a suitable solvent such as water, acetone, chloroform etc, in

which the polymer is soluble and the clay is dispersible. When the polymer

solution and the clay-dispersed solution are mixed, the polymer chains will be

intercalated between the clay layers by replacing the solvent molecules.

Intercalated polymers will remain in the clay layers upon the removal of solvent. It

is reported that the increase in entropy by the desorption of solvent molecules is

the driving force for the intercalation of polymer from solution [114, 171-175].

Water soluble polymers such as poly (ethylene oxide) [176], poly (vinyl alcohols)

[177], poly(ethylene vinyl alcohol) [178], polyvinyl- pyrrolidone [179] etc: have

been intercalated between the clay layers by this method. Nanocomposites based

on high-density polyethylene [180], polyimide [130,181], PCL [182], PLA [183]

etc have been synthesized by this method using non-aqueous solvents. The major

27


advantage of this method is that it offers the possibility to synthesize intercalated

nanocomposites based on polymers with low or even without polarity.

1.6.2 Properties of Polymer clay nanocomposites

Polymer layered silicate nanocomposites have attracted great interest, both

in industry and in academics, because they exhibit remarkable improvement in

material properties compared with virgin polymer or conventional micro and

macro composites [107-113]. Conventional composites usually require a high

content (>10%) of the inorganic fillers to impart the desired mechanical properties.

Such high filler levels increase their density of the product and can cause the

deterioration in properties through interfacial incompatibility between the filler and

the organic material. Besides, processability worsens as filler content increases. In

contrast, nanocomposites show enhanced thermo mechanical properties even with

a small amount of layered silicate (~5%). Improvements comprise higher modulus

[122-125], increased strength, heat resistance [126], decreased gas permeability

[127-131], reduced coefficient of thermal expansion [184,185] and decreased

flammability [132-139]. The main reason for this improved property in

nanocomposites is the large interfacial interaction between the matrix and layered

silicate and also the high aspect ratio of the dispersed clay particles. Polymer clay

nanocomposites have been studied with different polymer matrixes such as

polystyrene [ISO], poly (£-caprolactone) [162,163], poly (ethylene oxide) [176],

polyamide [131,142-146], polyimide [130,181], epoxy [186-192], polysiloxane

[193], and polyurethane [194].

1.6.2.1 Mechanical properties

Mechanical properties of polymer--elay nanocomposites depend on the

microstructure in which how the clay layers are dispersed in the polymer matrix.

Generally the well dispersion of the clay particles in the polymer matrix yields

enhanced tensile modulus, storage modulus and tensile strength. [195,196]. Even

though, the tensile strength and modulus tend to increase with increasing clay

content, the increasing trend is more noticeable for the tensile modulus. The

28


reinforcing effect of clay layers on the tensile modulus is mainly due to the high

modulus and high aspect ratio of the dispersed clay layers. This will provide large

interfacial interaction between clay layers and polymer matrix. Kojima et al

[131,197] reported 100% improvement in young's modulus in polyamide 6

nanocomposites with less than 5 wt% of clay particles. In PCL clay

nanocomposites, the young's modulus was increased to 270% with the

incorporation of 8 wt% organoclay [198]. Similarly Shelley et al [199] observed

175% increase in tensile strength along with 200% increase in tensile modulus

with nylon 6 nanocomposites containing 5 wt% of clay. Basara et al observed

17.2% increase in tensile modulus when 7 wt% of Cloisite 30B was incorporated

to epoxy matrix [200]. The study on the morphology and mechanical properties of

epoxy system with octadecylammonium ion-modified MMT showed an

improvement in modulus and fracture toughness (KId and exhibited a mixed

intercalated/exfoliated structure by Becker et al [201]. Pluart et al [202] also

reported similar improvements in the tensile strength and stiffness as well as

fracture toughness of DGEBA resin resulting from clay additions. Zhang et al

found 88% and 21 % improvement in impact strength and tensile strength

respectively for epoxy clay nanocomposites having 3 wt% organoclay [203].

Ingram et al [204] observed an increase in. storage modulus for DDS and DDM

cured DGEBA epoxy clay nanocomposites. Similar observations are made by

Hussain et al [205] for aromatic amine cured DGEBF epoxy clay nanocomposites.

The storage modulus and fracture toughness were found to increase with increase

In clay concentration for epoxy clay nanocomposites cured by

diethyltoluenediamine [206].

However, many authors [110,207-209] have reported a reduction in tensile

strength, impact strength, fracture toughness as well as the strain at break of the

polymer clay nanocomposites with the addition of clay particles. There are several

reasons to explain the decrease in properties as clay content increases .One reason

is the stress concentration effect of agglomerated clay particles at higher clay

loadings. The agglomeration of clay particles at higher clay concentration also

results lower in mechanical properties due to lowering of filler surface area and

29


lower polymer/clay surface interaction. Another reason is that as clay content

increases, the viscosity of the system increases resulting in heterogeneity and

nanovoids formation due to the entrapment of air bubbles during sample

preparation. Miyagawa et al [210,211] have observed a decrease in impact strength

with increase in clay content for anhydride and amine cured epoxy

nanocomposites.

1.6.2.2 Thermal properties

Polymer clay nanocomposites are known for its high thermal stability and

flame retardancy. The improved thermal stability is attributed to the action of clay

layers as superior insulator and mass transport barrier to the volatile products

generated during decomposition as well as assisting in the formation of char after

thermal decomposition [107, 212-214]. The slowing down of the escape of the

volatile products in nanocomposites is because of the labyrinth effect of the silicate

layers in the polymer matrix [215,216]. The improved thermal stability of the

polymer clay nanocomposites has been reported for various types of organoclays

and polymer matrices. Zhang et al [203] reported that the heat distortion

temperature and thermal decomposition were heightened from 124°C and 348°C to

133°C and 373°C respectively for epoxy clay nanocomposites with 5 wt% clay

compared to pristine epoxy matrix. Blumstein [217] observed a 40-50°C higher

decomposition temperature for PMMA- clay nanocomposites. The improved flame

retardancy of the polymer clay nanocomposites may be attributed to the formation

of thermal insulating and low permeable char residue of clay layers at the outer

surface of the nanocomposite during combustion and acts as a protective barrier by

reducing the heat and mass transfer between the flame and polymer. The char

residue of clay layers will also reduce the oxygen uptake and the escape of volatile

gases produced by polymer degradation. Camino et al [218] reported the

mechanism of the improved fire retardany of epoxy clay nanocomposites cured by

methyl tetrahydropthalein anhydride. This is due to the formation of protective

skin created by ablative reassembling of the clay layers as well as the chemical

structure of the clay.

30


Many investigators reported that the Tg of the polymer either increase or

decrease with the addition of clay particles. The increase in Tg is attributed to the

slower segmental motion due the polymer chains being anchored to the surface of

the clay. Lu et al [219] and Miyagawa et al [220] observed an increase in Tg with

the addition of clay particles to epoxy resin. Liu et al [206] observed a decrease in

Tg with increase in clay content. This was attributed to the plasticizer effect of the

clay modifier.

1.6.3.3 Barrier properties

Clay layers in the polymer matrix can act as an effective barrier to the

penetrants. The enhanced barrier property of polymer nanocomposites is due to the

labyrinth or tortuous path (Fig. 1.11) that retards the diffusion of gas molecules

through the polymer matrix. Polyimide clay nanocomposites exhibited reduction in

gas permeability with small fraction of clay layers [130].

Fig 1.11 Zigzag diffusion (tortuous) pathway of a gas through clay-based polymernanocomposites

Neilson's equation is proved to be a reliable estimate [221] of gas

permeability of polymer-layered silicate nanocomposite systems. The Neilson's

equation is as follows

where Pn represents the permeability of the resulting nanocomposite and Pm

represents the permeability of the matrix polymer, ~ is the volume fraction of clay

31


platelets, a = ~ , is the aspect ratio of the clay layers and Land t are the lengtht

and thickness of clay layers respectively. Due to the mismatch of the experimental

values with theoretical predictions, the tortuosity factor was corrected by the

modified Neilson's model to include the orientation factor as proposed by

Bharadwaj [128].

According to Bharadwaj, the permeability of a nanocomposite (Pn)

containing clay stacks of length (L) (L also being the length of the individual clay

platelets) and thickness (t) is related to the permeability ofthe pure matrix (Pm) by

where $ is the volume fraction of clay platelets, S is defined as an order

parameter (S=~(3cos2e-1» representing the orientation of the tactoids in the2

matrix and 8 represents the angle between the direction of preferred orientation (n)

and the sheet normal (p) unit vectors as shown in Fig. 1.12. The order parameter

(S) can range from 1 to -1/2. If S=I(8 =0), the tactoids will be aligned

perpendicular to the gas flow direction (in this case, the model is equivalent to

Neilson's model). When S = 0 (8 =57), the tactoid will be distributed randomly

and if S = -1/2 (8 = 1t/2), the tactoids will be aligned in the gas flow direction.

5=1

Fig 1.12 Order parameter of the clay layers

32

Introduction

1.7 Epoxy clay nanocomposites

Chapter I

Among polymer layered silicate nanocomposites, epoxy based systems have

been studied extensively by several researchers [187,222-231]. There are many

factors which influence the structure and properties of nanocomposites, for

instance nature of curing agent, nature of clay, curing conditions, processing

method etc. It was found that when anhydride was used as curing agent, an

exfoliated morphology while diamino diphenyl methane (DDM) gave an

intercalated morphology, since anhydride is a liquid and can easily diffuse into the

clay gallery unlike DDM which is a solid [222]. Similar observations were made

by Xu et al [223] for diethylenetriamine and tung oil anhydride. Kornman et al

[224] found that aliphatic amine cured epoxy produced an exfoliated morphology

compared to cyclo aliphatic amine cured epoxy because of the higher reactivity of

former. The diffusion rate and reactivity of the curing agent also influence the

exfoliation of clay. They also studied the influence of the nature of the clay on the

structure of epoxy clay nanocomposites. Organic modified montmorillonite clay

with a low CEC showed an exfoliated structure compared to clay with high CEC,

which showed an intercalated structure during the swelling of clay in the epoxy

resin for 24 hours prior to curing. This is due to the low amount of organic

modifier (octadecyl ammonium chloride) present in the former clay provides more

space for DGEBA molecules. So the self polymerization of epoxy can occur in

large extent and causes the diffusion of new DGEBA molecule between the clay

layers leading to the exfoliation of clay [225].

Intercalated nanocomposite is generally produced with quaternary and

tertiary alkyl ammonium surfactants due to low bronsted acidity of the surfactants

[187]. The fixed layer separation of clay layers is unable to provide optimum level

of reinforcement. Exfoliated nanocomposite is produced with primary and

secondary alkyl ammonium surfactants or quaternary surfactants containing

hydroxyl groups due to the high bronsted acidity of the surfactants [187]. In

exfoliated nanocomposites, the clay layers will be sufficiently separated and

randomly oriented to allow full interfacial bonding with matrix resin to improve

the properties of nanocomposites. Similar observation was made by Lan et al [187]

33


in the case of m- phenylene diamine DGEBA epoxy resin nanocomposites

containing 5% alkyl ammonium modified clay. They reported that in the case of

primary and secondary ammonium modified clay, the rate of intragallery

polymerization occurs at a faster rate compared to extragallery polymerization

leading to an exfoliated structure while the extragallery polymerization is in faster

rate in the case of tertiary and quaternary ammonium modified clay results an

intercalated morphology. It is also reported that as the chain length of alkyl

ammonium (clay modifier) increases, the morphology changes from intercalation

to exfoliation morphology [187]. Similarly Wang et al [232] also reported that the

curing speed of the inter and extralayer epoxy amine reaction is the key factor for

the synthesis of exfoliated epoxy clay nanocomposites. Kong et al [233] studied

the exfoliation behaviour of epoxy clay nanocomposites varying the electro

negativities of the aromatic diamine curing agent and curing temperatures. Epoxy

clay nanocomposites based on DGEBA/PDA and DGEBA/MDA systems

produced an intercalated structure due to the high reactivity of the curing agents

which induces faster gelation in the extragallery region, while DGEBA/DDS

system gave exfoliated structure due to the low reactivity of the curing agent

thereby slowed down the extragalley gelation and provides enough time for

intragallery polymerization.

The processing methods can also influence the clay morphology. The usual

processing methods to disperse the clay layers in epoxy matrix are mechanical

stirring, ultrasound sonication [226-227], high shear mixing [228-230], ball milling

[231] etc. Lam et al [226] reported that 10 minutes ultrasonication result an

optimum micro hardness at 4 wt % clay containing epoxy nanocomposite. Zunjarro

et al [227] reported that high speed shear mixing yielded better mechanical

properties compared to ultrasonication, even though both methods gave exfoliated

morphology. Yasmin et al [228] found that epoxy clay nanocomposites processed

by three roll mill is efficient in achieving high levels of exfoliation and dispersion

of clay particles within a short period of time. The modulus of the nanocomposites

found to increase with increase in clay concentration. However, the tensile strength

was decreased with addition of clay particles compared to pure epoxy due to the

34


occasional occurrence of nano to micro sized voids in the microstructure. They

observed that effective degassing during processing will enhance the tensile

strength of the resultant nanocomposites. Chen et al [234] prepared a fully

exfoliated layered silicate epoxy nanocomposites by the combination of high shear

mixing and ultrasonication in the presence of acetone. Lu et al [230] observed a

decrease in impact strength and flexural strength for 4,4'-diamino diphenyl sulfone

(DDS) cured epoxy clay nanocomposites processed by mechanical stirrer as well

as high speed emulsifying and homogenizing mixer (HEHM) where as processing

by HEHM followed by ball milling improved the mechanical properties. It was

reported that epoxy-DOS-clay nanocomposites processed by high pressure mixing

method showed dramatic increase in fracture toughness compared to direct mixing

method [235]. However, the glass transition temperature decreases as the clay

content increases.

Epoxy resin is found to be an attractive polymer matrix for cryo tank

applications due to their low cure shrinkage. However, microcracking of the matrix

due to the mismatch in coefficient of thermal expansion between fibre and matrix

of the polymer matrix composites remain unresolved for the use in cryo tank. The

microcrack can facilitate in the permeation of cryogen through the inner skin of the

tank to the honeycomb core. In the case of X-33, microcracking resulted in the

catastrophic destruction of the tank [236]. Epoxy clay nanocomposite is a

promising and ideal substitute to overcome both the mismatch in CTE and

permeability. NASA Glenn Research Center has developed the epoxy-clay

nanocomposite with up to 70% lower in hydrogen permeability and 25% decrease

in the CTE than that of the neat epoxy resin [185]. Zilg et al [237] found higher

toughness for an intercalated structure and higher modulus (stiffness) for

completely exfoliated structure for nanocomposites based on anhydride cured

DGEBA epoxy. Qi et al [238] reported 25% improvement in fracture toughness

with the incorporation of 5% Cloisite 30B in epoxy matrix.

Chin et al [239] reported that an exfoliated nanocomposites is formed when

the DGEBA epoxy with octadecyl ammonium modified MMT is cured with less

35


than stoichiometric amount of meta-phenylene diamine (MPDA) or with

autopolymerisation without curing agent. On other hand an intercalated

nanostructure is formed when it is cured with equimolar of higher concentration of

MPDA. Extragallery cross linking has been dominated in the case of higher

concentration of curing agent, resulting in intercalated nanocomposites. Liu et al

studied [240,241] the effect of mixing method to improve the dispersion of clay in

epoxy and observed a significant improvement in fracture toughness at 1 wt% of

clay loading for tetraglycidyl diamino diphenyl methane (TGDDM) epoxy-DDS

clay nanocomposites synthesized with high pressure mixing method compared to

direct mixing method. The Tg of the nanocomposites is found to decrease with

increase in clay content. The effect of temperature, speed and time at the pre

mixing step during dispersion on intercalation and exfoliation of clay in the epoxy

resin have been studied by Ngo et al [242]. Even though the above premixing

parameters have not any significant effect on the intercalation of organoclay at the

pre-mixing step, they have a positive effect on the intercalation/exfoliation of

nanoclay at the curing step. Pluart et al [243] also demonstrated that adequate

compatibilisation of montmorillonite followed by swelling of clay galleries by

epoxy was necessary to obtain intercalated/exfoliated nanocomposites. In another

study [202], they correlated the morphology and mechanical properties of epoxy

clay nanocomposites and observed an improvement in stiffness for exfoliated

nanocomposites. However, a very interesting stiffness/toughness balance was

shown by the intercalated nanostructures without lowering their Tg.

Wang et al [244,245] developed a new method, namely slurry

compounding for the preparation of highly exfoliated epoxy clay nanocomposites

using pristine clay, which involved the transfer of clay water suspension to epoxy

resin by solvent exchange step and silane modification step. The important feature

of the technique was that very little amount of organic modifiers was enough to

facilitate the exfoliation of clay, in contrast to conventional organoclays, which

normally contains at least 25-45% of organic surfactants. The resultant

nanocomposites showed improvement in fracture toughness, young's modulus,

storage modulus and Tg. The formation of a large number of microcracks and the

36


increase in fracture surface area due to crack deflection are the major toughening

mechanism in the nanocomposites. The influence of the change of the functionality

of the epoxy compounds and the amine curing agents and the effect of the variation

of the concentration of the clay on the physical properties of epoxy clay

nanocomposites were reported by McIntyre et al [246]. The storage modulus and

Tg of the nanocomposites were increased with the addition of Cloisite 30B for

DGEBA epoxy cured by tryethylenetetramine (TETA), DDS and DDM while a

lowering of Tg was observed for TGDDM/DDS system. Among the DGEBA

system, DDS cured system showed improved Tg and modulus compared to other

curing agents. All composites haven't shown any significant improvement in the

thermal stability by the addition of clay particles.

Kaya et al [247,248] reported that the Tg, storage and loss moduli of the

neat epoxy were increased by the incorporation of clay particles. The incorporation

of unmodified clay (MMT) into the epoxy resin didn't affect the Tg value, while

the addition of 3 wt% of organically modified clay (OMMT) increased the Tg by

about 15°C due to the better exfoliation of clay in the epoxy matrix. It is also

observed that the epoxy clay nanocomposites containing OMMT clay particles

exhibited better optical transparency than those with MMT. Flame retardancy of

the epoxy was increased by the addition of clay particles and the burning rate is

decreased by 38 and 58% for MMT and OMMT nanocomposites respectively for

10% clay loading.

Influence of clay surface modification on the structure of the epoxy clay

nanocomposites using a hardener of polyoxypropylenediamine have been

investigated by Ryznarova et al [249]. They observed that the difference between

the curing rates inside and outside clay gallery was crucial for achieving

intercalation/exfoliation of the prepared nanocomposites. Protonated and

functionalized clay modifiers catalyzed the intragallery polymerization of epoxy

resulted in the reduction in gelation times, gradual increase in d spacing during

curing and high degree of dispersion leading to increased elongation at break and

toughness. In contrast, non functionalized alkyl ammonium ions were umble to

37


catalyze the intragallery polymerization resulted in the faster extragallery

polymerization leading to only partially intercalated nanocomposites.

Triantafyllidis et al [250] reported that the epoxy clay nanocomposites prepared by

the incorporation of homoionic organic clays exchanged with relatively short chain

di- or triamines and mixed-ion organic/inorganic clays partially exchanged (35%)

with long chain diamines modified by di- or triamines resulted in intercalated

structures with improved young's modulus and storage modulus. On the other

hand, homoionic organic clays exchanged with long chain diamines and triamines

resulted in exfoliated nanocomposites, but with compromised mechanical

properties especially reduced Tg due to the plasticizing effect of the long chain

amine modifiers. Wang et al [251] investigated the effect of clay concentration on

the morphology and properties of epoxy clay nanocomposites prepared by in situ

polymerization under mechanical stirring followed by ultrasonication and observed

a decrease in layer space along with aggregation of clay particles with increase in

clay concentration due to the increase in viscosity. Thermal decomposition

temperature remained unchanged and the Tg decreased, whereas the storage

modulus increased with increase in clay concentration. According to them, the

improved storage modulus is due to the stiff filler reinforcement with partial

exfoliation.

The durability studies of epoxy montmorillonite clay nanocomposites under

room temperature, in hot and cold conditions showed that the mechanical

properties were found to decrease with increase in time. 2wt% clay filled epoxy

nanocomposites showed enhancement in properties with relatively less number of

cracks and better interfacial bonding in all conditions over its neat counterpart

[252]. Lakshmi et al [253] investigated the thermal stability and structural

characteristics of different epoxy clay nanocomposites using hexadecyl ammonium

and phosphonium clay and DDS as hardener. The ammonium modified clay epoxy

system showed appreciable mechanical and glass transition temperature properties

while phosphonium modified clay epoxy system exhibited highest thermal

resistance properties compared with unmodified epoxy systems. Similarly, the

38


higher thermal resistance properties of the alkyl phosphonium clay epoxy

nanocomposites were reported by Wei et al [254].

1.8 Polyimide clay nanocomposites

Polyimides (PI) find applications in many fields such as electronic, optical,

automotive and aerospace industries, because of their excellent thermal stability,

high mechanical performance and outstanding electrical performance [255]. The

properties of the polyimide can be further improved by the incorporation of clay

particles. The performance of PI clay depends on the clay dispersion, the type of

alkyl ammonium ions used as an organic modifier, the silicates present in the clay,

and the PI type [130, 256-267]. There have been a number of attempts to improve

the various properties of PI/clay hybrids, such as barrier performance in terms of

the tortuous path of penetration of gas molecules [256,257], thermal expansion

[130,256], mechanical and thermal properties [258,259,268], and imidization

reaction kinetics' time and chemical resistance [259,269]. PI/clay hybrids have

also been reported to have a interiayer spacing of 13.2A° independent of the chain

length of the intercalated alkyl ammonium cation after thermal imidization at

around 300°C, although the interiayer spacing of uncured poly(amic acid) (PAA)

hybrids with organically modified clays shows some variations in interiayer

spacing [109, 110, 257]. Lan et al [257] concluded that the change in interiayer

spacing was related not to the thermal degradation of onium ions, but to the

eviction of onium ions from the clay interiayers, and that the characteristic spacing

of 13.2A° was responsible for the intercalation of PI molecules in a flattened

conformation.

Yano et al [130] have studied PI-clay hybrid based on pyromellitic

dianhydride (PMDA) and 4, 4' -diaminophenyl ether in dimethylacetamide. They

found that a several-fold reduction in permeability of small gases was achieved,

despite having an intercalated morphology. A parallel study by Lan et al. [270]

confirmed the findings ofYano et al [130]. Chang et al [271] reported a study ofa

thermoplastic PI nanocomposite based on PMDA and benzidine. They prepared

nanocomposites using different organically modified montmorillonite and achieved

39


significant improvements in tensile and barrier properties. Yano et al reported that

2 wt % addition of the montmorillonite (MMT) in PI, the permeability coefficients

of H2, O2, and water vapor have been reduced to less than half of these of the

pristine PI, and the CTE is lowered to 90% of the neat PI [130].

The commonly used organo-modification agents are long carbon-chain

alkyl ammonium salts. It has been widely accepted that the interlayer spacing of

organo-Iayered silicates depends greatly on the length of the carbon chain [123].

Although these modification agents have been gaining significant success in the

preparation of polymer/MMT nanocomposites, their common shortcoming is the

poor thermal stability. Xie et al [272,273] have studied the thermal stability of

MMT modified by long carbon-chain alkyl quaternary ammonium ions using

TGA-MS and found that the on-set decomposition temperature of the resultant

OLSs was approximately 180°C. Unfortunately, the preparation and processing of

most of the polymer/OLS nanocomposites require a temperature much higher than

this value, and the thermal decomposition of the long carbon-chain alkyl

quaternary ammonium salts is inevitable. Delozier et al [274] observed that during

the preparation of polyimide/clay nanocomposites, the decomposition of the

organic modifier led to the 'collapse of the clay particles into larger agglomerates'.

This may affect the morphological structure, properties and service life of

nanocomposites. Therefore, the thermal stability of organic modifier has a

significant effect on the preparation, performance and application of

nanocomposites.

1.9 Epoxy syntactic foam clay nanocomposites

Hollow particle (Microballoon) filled composites known as syntactic

foams, are finding increasing use in aerospace, marine and defense applications

due to its low density, high compressive strength, low CTE and high resistance to

moisture absorption [274-278]. Most extensively used hollow particles are glass

microballoons because of its high specific strength and energy absorption

properties [279]. The matrix resin usually preferred for the manufacturing of

syntactic foam is epoxy resin due to its high strength and stiffness, thermal stability

40


and creep resistance [51, 53]. However, their brittle nature due to its high cross

link density and poor resistance to crack propagation limits the applications in high

perfonnance structures. Due to the increased demand of syntactic foam in

sandwich structures, there is a need to improve the modulus and fracture toughness

properties within the low-density region.

Wouterson et al [280-282] reported that creation of different

microstructures by using different types of microspheres and inclusion of short

microfibres are the effective methods to improve the fracture toughness of the

syntactic foams. They observed an optimum fracture toughness with 20-30 vol%

microspheres and thereafter it decreases and they also found 95% increase in

fracture toughness by the addition of 3 wt% short carbon fibres to 30 vol%

microsphere filled epoxy syntactic foam. But Kim et al [283-284] found a

reduction in fracture toughness and flexural strength of the syntactic foam with

increase in the volume fraction of the glass microballoons, accompanied by

improvement in impact properties. It is also known that modulus, toughness and

thermo mechanical properties of epoxy resin can be significantly enhanced by the

addition of nanoclay [107-112]. Gupta et al [285] observed 80 and 125% increase

in toughness, measured as area under the stress strain curves, for 2 and 5%

nanoclay respectively in syntactic foam containing microballoons of density

220kg/cm3• However, they observed a decrease in compressive strength by 10-20%

by the incorporation of 2 vol % nanoclay, whereas 5% nanoclay has given nearly

same level of strength. Later, Wourtson et al [281] reported that the micro fibre

toughening is more effective than toughening with nanoclay except 1 wt%

nanoclay, which showed equal, or better fracture properties compared to short fibre

reinforced syntactic foam.

Maharsia et al [286] has studied the effect of rubber particles and nanoclay

particles on the flexural properties of epoxy syntactic foams. The flexural strength

and fracture strain of the syntactic foam are reduced by 11 % and 15% by the

addition of 2% nanoclay irrespective of the density of microballoons, while 22%

increase in flexural strength is shown by the incorporation of 5 % nanoclay to the

41


low density syntactic foam. But they observed an increase in flexural strength and

fracture strain with the addition of rubber particles to epoxy syntactic foam. It has

been reported that the crack propagation resistance and fracture toughness of the

epoxy glass microballoon system can be improved with the incorporation of

reactive liquid rubber [287]. Gupta et al [288] reported a significant increase in

compressive toughness and energy absorption for the crumb rubber particles

(particles obtained from waste rubber) incorporated epoxy syntactic foam.

However, they observed 50% decrease in young's modulus and 10% reduction in

compressive strength. An enhancement in impact energy of both initiation and

propagation are observed by the incorporation of crumb rubber, microfibre and

nanoclay to the epoxy syntactic foam, due to the toughening effect of crumb

rubber, strengthening effect and stress distribution of microfiber and nanoclay

[289-291]. Woldesenbet [292] reported that the incorporation of nanoclay in

syntactic foam has enhanced the maximum load and impact initiation energy of

syntactic foams with higher (K46) and lower (S22) shell thickness.

1.10 Carbon fibre reinforced epoxy composites

Carbon fibre reinforced plastics (CFRP) have been emerged as an ideal

substitute for metals in many structural applications. High specific fracture

toughness and strength and good dimensional stability of these composites make

them more attractive in aerospace and automobile industries [293-297]. Epoxy

resin is the commonly used matrix resin for the development of CFRP, because of

their excellent adhesion, strength, low shrinkage, corrosion resistance and

processing versatility [51,53].

NASA Glenn Research Center has developed the epoxy-clay

nanocomposites with up to 70% lower in hydrogen permeability and 25% decrease

in the CTE than that of the neat epoxy resin. Filament-wound carbon-fiber

reinforced tanks made with this nanocomposite have showed fivefold lower helium

leak rate than the corresponding tanks made without clay [185]. Chowdhury et al

[298] reported that the mechanical properties as well as the thermal properties were

improved for the nanoclay filled carbon fibre reinforced composites having

42


optimum concentration of 2 wt%, but the properties were declined in the case of 3

wt% nanoclay loading. This is due to the fact that at higher concentration of clay

loading, nanoclays start to agglomerate and act as flaws and crack initiation sites.

Gilbert et al [299,300] and Timmerman et al [301] observed an increase in the

fracture toughness and mechanical properties with the addition of metal and

inorganic particles to the epoxy based carbon fibre reinforced composites system.

Mohan et al [302] observed an increase of 12% in tensile modulus and 8% in

tensile strength when 1.5 wt% of alumina nanoparticles were dispersed in S-2 glass

composites. An improvement in interlaminar shear strength was observed by

Miyagawa et al [303] in carbon fibre reinforced composites by using biobased

epoxy clay nanocomposites as matrix. However, there was no change in the

flexural strength and modulus with the incorporation of nanoclay. Mode I fracture

toughness and Gte of DGEBA/ diethyltoluene diamine/unidirectional carbon fibre

composites was improved by 25% and 50% respectively with the addition of 7.5%

of clay loading compared to neat system. Investigation on the ternary carbon fibre

epoxy nanocomposite based on a DGEBF epoxy resin/clay showed a minor

improvement in flexural strength at low clay concentrations but as clay content

increases the flexural strength decreases. This is due to the formation of voids in

the matrix at higher clay concentration.

1.11 Scope and Objective of the work

The toughening effect of thermoplastic and nanoclay separately on the

thermo mechanical properties of epoxy systems has been studied enormously by

several researchers. Blending epoxy with thermoplastic such as Hydroxyl

terminated polyether ether ketone with pendant methyl groups (PEEKMOH),

Hydroxyl terminated polyether ether ketone with pendant tertiary butyl groups

(PEEKTOH), Polyether sulfone (PES) etc is an effective method to improve the

fracture toughness without sacrificing other thermo mechanical properties [25-31].

On the other hand, clay reinforced epoxies have shown improved modulus, flame

resistance and thermal stability as well as reduced gas permeability and coefficient

of thermal expansion [107-113]. However no evidence of drastic increase in

43


fracture toughness was reported in the open literature for epoxy clay

nanocomposites. On the other hand more than 100% increase in fracture toughness

was observed for the thermoplastic toughened epoxy systems [25-31]. Therefore, if

clay particles are dispersed in a thermoplastic toughened epoxy, the resultant

ternary nanocomposites have a synergistic improvement in toughness, modulus,

coefficient of thermal expansion and gas barrier properties.

Frohlich et al [304] reported an improved toughness in rubber toughened

hybrid epoxy nanocomposite compared with nanocomposite without rubber. The

Tg was lowered drastically due to the presence of rubber. Balakrishnan et al [305]

observed that the ductility of epoxy resin was enhanced without considerable

reduction in the modulus and strength when organoclay and rubber dispersants

were added to epoxy resin. Isik et al [306] studied the mechanical properties of

epoxy-polyetherpolyol-organically treated montmorillonite nanocomposite with

respect to polyetherpolyol content and clay content. They observed an increase in

Tg and young's modulus with respect to clay content. Peng et al [307] investigated

the impact of clay on the phase morphology of epoxy/poly (ether imide) ternary

hybrid nanocomposites.

Hence, the present study focuses on the incorporation of clays into high

performance thermoplastic (PEEKMOH, PES, PEEKTOH) toughened epoxy,

which provides synergistic effect on toughness, modulus, coefficient of thermal

expansion and gas barrier properties. The surface morphology and thermo

mechanical properties of the above ternary nanocomposites are evaluated with

respect to the nature and concentration of both thermoplastic and clay. The present

work also emphasises on the preparation and property evaluation of syntactic foam

and unidirectional carbon fiber composites from the thermoplastic toughened

epoxy clay systems. The study further highlights the synthesis of high temperature

resistant phosphonium modified clay and the effect of clay and thermoplastic on

the thermo mechanical and barrier properties of the polyimide system.

Applications of these materials In the space industry include

nanostructured, heat-insulating layers for rocket engines, cryogenic storage tanks

to restrict the leak of cryogens and also in the ablative environments in the rocket

44


motors (owing to their increased thermal stabilities). Thus thermoplastic toughened

nanocomposites offer potential in the development of lightweight, durable LH2

tankage, Helium gas bottles, high performance satellite structures for a wide range

of applications from reusable launch vehicles to fuel cell powered aircrafts.

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