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PET blends and glass fibrecomposites

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LOUGHBOROUGH UNIVERSITY OF TECHNOLOGY

LIBRARY AUTHOR/FILING TITLE

{<061 N SO,..J It M ---------- ----- ---- -----j-------------- ----.---

ACCESSION/COPY NO.

VOL. NO. CLASS MARK

- 6 JUL 199U 26 JUN 199B

ii 5 JUl 1

- 5 JUL 1~91_ 8 0 .T 1Q!l~

i 1 ,Cl 1 - I t>t

I

PET BLENDS AND GLASS FIBRE COMPOSITES

by

A.M. ROBINSON

A Masters Thesis Submitted in Partial Fulfilment of the

Requirements for the Award of MASTER OF PHILOSOPHY

of the

Loughborough University of Technology

February 1986

Supervisors: Professor A.W. Birley, Institute of Polymer Technology

Mr. B. Haworth, Institute of Polymer Technology

t) by A.M. Robinson (1986)

---'-"--L.\AllJhll>., u .... !~ Vnh,ei-'"

.t t lIIoc..t,flt>1091 LIt.. .. J .... 01AM.L I &1::> f:1_ I "c<.

0102..7:>7/01 ....

Dedicated to my loved ones; Gill and my parents.

ACKNOWLEDGEMENTS

I would like to express my sincere gratitude to Professor A.W. Birley

for the opportunity to undertake this programme. His enthusiasm, knowledge

and tireless effort was a continual inspiration.

I am also indebted to Mr B. Haworth who could always be relied upon

to encourage and for the advice that arose from our numerous discussions,

which was invaluable.

The completion of this programme was greatly aided by Mr P. Ramsey

and Mr. A.J. Davis who were always available to steer me in the right

direction with all technical problems. I would like to thank Mr R. Owens

for his time, . knowledge and friendship.

A special thank you to Judi the typist for an excellent job cheerM

fully done.

Finally, I would like to thank Ford Motor Company for the financial

support without which this programme would not have been possible.

A.M.R.

ABSTRACT

Polyethylene terephthalate (PET) is a potentially valuable

material which can be reclaimed and recycled into a variety of high

value applications, such as in the automotive industry. The physical

properties required for applications of PET can be attained by blend­

ing and/or addition of reinforcing fibres. The literature for blends

and composites of PET is quite extensive, however there are areas that

have not been investigated. It is necessary to explore the effect of

reinforcing a blend of PET and bisphenol-A-polycarbonate (PC) with

glass fibres for example.

Recycled PET has been shown after thorough drying, to have very

similar properties to that of virgin PET. This means that reclaiming

and recycling of PET is a viable propOSition. With care in processing

the reduction in molecular weight (degradation) can be ignored.

The composites of PET and glass fibres have been shown to exhibit

exceptional flexural strength accompanied by good impact""" strength.

When PC is added to this composite in the proportion 20% by weight the

impact strength increases dramatically.

A wide variety of nucleating agents are suggested as being effective

for PET; in practice the choice is limited basically to salts of mono­

carboxylic or polycarboxylic acids. These gave a high degree of

crystallinity at high nucleation rates both of which are very important

in injection moulding of PET and for its ultimate physical properties.

1. 1 •

1.2.

1.3.

CONTENTS

AIM

CHAPTER ONE

-LITERATURE REVIEW

INTRODUCTION

RAW MATERIALS

1.2.1. POLYETHYLENE TEREPHTHALATE

1.2.1.1. MOLECULAR WEIGHT

1.2.1.2. CONTAMINANTS

1.2.1.3. WATER

1.2.2. GLASS FIBRES

, 1.2.2.1. FIBRE CONTENT ,

1.2.2.2. FIBRE LENGTH

1.2.2.3. COUPLING

1.2.3. IMPACT MODIFIERS

1.2.4. CRYSTALLISATION NUCLEANTS

1 .2.5. POLYCARBONATE

1.2.5.1. MOLECULAR WEIGHT

PROCESSING

1.3.1. DRYING

1.3.2. PRECOMPOUNDING

1.3.3. INJECTION MOULDING

1.3.3.1. MOULD TEMPERATURE

1.3.3.2. GATE DESIGN

1.3.3.3. INJECTION SPEED

Page No.

5

5

8

8

10

12

15

15

19

20

21

25

21

29

30

30

30

32

33

34

1 .3.4. COMPOUNDING-ADDITION OF FIBRES 35

1.3.5. PROCESSING EFFECTS ON PROPERTIES 36

1.3.5.1. CRYSTALLISATION 36

1.3.5.2. FIBRE BREAKAGE/ORIENTATION 41

1.3.5.3. BLEND CHARACTERISTICS 46

1.3.5.4. RECYCLING 48

1.4. GLASS FIBRE REINFORCEMENT 50

1.4.1. POLYETHYLENE TEREPHTHALATE REINFORCED WITH

GLASS FIBRES 50

1.4.1.1. STRENGTH 51

1.4.1.2. TOUGHNESS 74

1.4.1.3. STIFFNESS 84

1.4.1.4. ENVIRONMENT 88

1.4.2. PET/PC BLENDS I

106

1.4.2.1. STRENGTH 107

1.4.2.2. TOUGHNESS 110

1.4.2.3. STIFFNESS 110

1.4.2.4. ENVIRONMENTAL 111

2.1.

2.2.

CHAPTER TWO

CHARACTERISATION OF RAW MATERIALS

THE RECYCLE BOTTLE

2.1.1.

2.1.2.

2.1.3.

HYDROLYTIC DEGRADATION OF PET:EFFECT OF

RELATIVE HUMIDITY ON THE MEASUREMENT OF THE

MELT FLOW INDEX OF SCRAP PET

2.1.1.1. INTRODUCTION

2.1.1.2. EXPERIMENTAL

2.1.1.3. RESULTS

MEASUREMENT OF THE INTRINSIC VISCOSITY

OF THE MFI EXTRUDATE

2.1.2.1. INTRODUCTION

2.1.2.2. PROCEDURE

2.1.2.3. RESULTS

COMPARISON OF MFI AND I.V RESULTS

EFFECT OF DRYING TIME ON THE MFI OF SCRAP PET

2.2.1.

2.2.2.

2.2.3:--

DRYING PET REGRIND IN A DESSICATOR AT 23°C

DRYING IN AN OVEN AT 120°C

COMPARISON OF DRYING METHODS

2.3. STUDY OF CRYSTALLISATION UNDER VARIOUS IMPOSED

TEMPERATURE PROGRAMMES AND INVESTIGATING THE EFFECT

OF NUCLEATING AGENTS

112

112

116

116

116

118

119

119

120

120

121

124

124

124

127

127

3.1.

3.2.

4.1.

4.2.

4.3.

4.4.

4.5.

COMPOUNDING

CHAPTER THREE

PROCESSING

3.1.1.

3.1. 2.

3.1. 3.

3.1. 4

INTRODUCTION

SCREW CONFIGURATION

OPERATING PROCEDURE

COMPOUNDS

INJECTION MOULDING

CHAPTER FOUR

TEST METHODS

FLEXURAL TESTING

IMPACT TESTING

ENVIRONMENTAL STRESS CRACKING

ANALYSIS OF FRACTURE SURFACES

FIBRE LENGTH

132

132

133

133

134

135

136

137

137

138

140

140

140

5.1.

5.2.

5.3.

CHAPTER FIVE

EXPERIMENTAL RESULTS AND

FLEXURAL PROPERTIES

5.1.1. MAXIMUM FLEXURAL STRESS

5.1 .2. DEFLECTION AT BREAK

5.1 .3. FLEXURAL MODULUS

5.1 .4. FIBRE STRAIN

IMPACT PROPERTIES

ENVIRONMENTAL STRESS CRACKING

DISCUSSION

5.4. EXAMINATION OF FRACTURE SURFACE USING THE SCANNING

ELECTRON MICROSCOPE

5.4.1.

5.4.2.

PET + 30 w/w% GLASS FIBRES

PET + 50 w/w% GLASS FIBRES

5.5 FIBRE-LENGTH-DISTRIBUTION

REFERENCES

CHAPTER SIX

CONCLUSIONS

141

141

141

142

143

144

146

149

150

150

153

157

159

161

AIM

To recycle scrap PET from carbonated drinks PET bottles, blend with

bisphenol-A-polycarbonate and reinforce with glass fibre to produce a

composite/blend that will be of use in the automotive industry.

1. LITERATURE REVIEW

1.1. INTRODUCTION

Polyethylene terephthalate is a potentially valua':lle material which

can be reclaimed and recycled into a variety of high value applications,

such as in the automotive industry. It is estimated that forty-five

thousand tonnes of PET bottles were recycled in the United States of

America in 1984. Smaller scale recycling schemes are in operation in the

United Kingdom and Belgium.

PET is mainly used as a container for carbonated soft drinks, but it

is also used for distilled spirits, cosmetics and food. This container

has achieved rapid market acceptance and is now a familiar sight on super­

market shelves around the world.

The bottles used in this project were supplied by Carters Packaging

Limited, Long Eaton, produced from Eastman Kodak Kodapak (PET) 7352 (clear).

The physical properties of this grade are shown in Table 1.

PET is used pure in contained manufacture - few additive are required.

Of all plastic materials PET has a good range of barrier properties - it

keeps in carbon dioxide and keeps out oxygen and contaminants.

Table 1

Typical Physical Properties of

KODAPAKi&l PET 7352

Property Uni ts Test Method

Density, g/cm' ASTM D 1505

Bulk Density kg/m'

Poured

Vibrated

Melt Density, @ 285°C, g/cm'

Molecular Number average (Mn)

Molecular Weight average (Mw)

Intrinsic ViscoSity,Jl(g

Crystallinity, %

Crystalline Peak Melting POint,OC

Heat of Fusion, <- kJ /kg

Thermal conductivity, <W/mK<

Specific heat, kJ/kg

D 1895

DSC

C177 <

Value

1.4

785

850

1.2

23,000

46,000

0.74

50

245

59

0.25

80°C 1 .42

D 2766 1. 51

1.88

2.05

Acetaldehyde, ppm 3

Pellets:

Size and shape, mm 2.5 mm cube

KODAPAK is a trademark of Eastman Kodak Company

2

The PET used for carbonated drinks bottles such as Kodapak 7352 and

ICI Melinar B90 1847 (clear) are not coated. Bottles are sometimes coated

with PVdC latex if they are being used for sensitive products. The base

cup is high density polyethylene.

As automotive manufacturers seek solutions to weight reduction of

structural parts, reinforced plastics will see increasing application.

Thermoplastic matrix composites are of growing importance in the auto­

motive industries due to their ease of processing.

Du Pont produce a glass fibre reinforced grade of injection mouldable

polyethylene terephthalate under the trade name Rynite. Rynite has been

used for advanced technical applications in the automotive industry.

Rynite composites have an outstanding combination of stiffness and tough­

ness. Other features include a heat deflection temperature of 220°C at

1.8MPa, the same excellent chemical resistance (in reality hydrolysis is

a limitation), . electrical properties, and dimensional stability as other

PET grades and good processibility.

Polycarbonates have become firmly established as engineering plastics;

their main applications are in electronics, electrical engineering and

lighting engineering. Their use in automotive applications is hampered

by their poor petrol resistance, low temperature behaviour and hydrolysis.

This is unfortunate as certain properties of PC such as impact resistance

would be extremely useful in these applications. The limitations of PC

are overcome by blending them with other polymers (1) in this case PET.

The reinforcing of a blend with glass fibre is a very complex

situation. It is thought however that this blend/composite will provide

a very useful material for use by the automotive industry.

3

The applications for this material are mainly semi-structural, such as

spoilers and possibly body panels.

The attraction of recycling PET is obvious. By 1990 it is predicted

(1 a) that in the UK alone over 1, 000 million PET bottles will be produced

annually (present consumption is 40% of this). The cost of the discarded

bottles in Western Europe is £70 million annually.

Due to the complexity of the PET bottle, a special recycling system

is required. Such a system for PET recovery has been developed jointly

by Amberger Kacilinwerke GmbfL of West Germany and PlaslTech Limited,

Birmingham. The PET recovered is not suitable for food applications, but

is 99.8% pure although it should be remembered that any impurities can

cause degradation. However, it is claimed that the recovered PET has a

loss in intrinsic viscosity of <0.02, which is extremely low, considering

the processes involved and the opportunities for moisture uptake. If the

recycling is as successful as it is claimed, then the recovered PET is

obviously suitable for engineering applications. Recycled PET could

compare cost effectively with other engineering materials for purposes

that have not been considered to date due to its high cost. This can be

seen when comparing prices, the price of recycled PET is £500 per tonne

as opposed to virgin which is £1200 per tonne (1b).

4

1.2 RAW MATERIALS

1.2.1. POLYETHYLENE TEREPHTHALATE

The simplest route to PET formation is the esterification of

terephthalic acid with ethylene glycol which forms the monomer

(bis-' -hydroxy~thyl, terephthalate) which is polycondensed to give PET.

PET compounds are characterised by outstanding mechanical, thermal

and electrical properties and are ideal for applications demanding:-

High stiffness and surface hardness at high temperatures

Low tendency to "creep"

Low water absorption

Excellent dimensional stability

Good chemical and solvent resistance

Reduced flammability (specialised grades)

Typical properties of PET using Beetle PET as an example are shown

in Table 2. PET with these properties offers engineers new opportunities

for component design in areas such as the electrical, electronic, auto­

motive and domestic appliance industries. PET's resistance to oil and

grease and its excellent heat resistance make it particularly suited, to

automotive applications (see Table 3).

5

TABLE 2: TYPICAL PROPERTIES:' OF 'BEETLE' PE:!

PROPERTY

Nominal glass content

Nominal filler content

Special characteristics

Specific gravity

Mould shrinkage

Tensile strength

Tensile modulus

Elongation

flexural strength

Flexural modulus

Charpy icpac: strength Notched Unnotched

Heat distortion te::lperature @ 1.8MPa

Electric strength

Surface resistiVity

Volume resistivity

Comparative tracking index

Oxygen index

Flammability rating

3mm

1.5mm

Glow-wire rating

3mm

1.5mrn

). for natural materials

Test Method

ASTM D792

ISO R527

ISO R527

ISO R527

150"178

ISO 178

ASD1 0648

lEC 243

IEC 93

lEG 93

lEG 112

ASTH 2863-14

UL94

UL94

Unit

%

• -

%

MPa

GPa

%

MPa

GPa

• kJ/m kJ/m 2

'c

MV/m

log109hmcm

109,OohmClll

%

6

Unfilled grade

PET 100

Amorj:lhous

1.34

0.2-0.5

50

2.5

16

80

2.5

8 no break

65

13

15

15

210

28

H8

750

650

TABLE 3 Resistance to cheoicals and oils (30days)

Engine oil

Turbine oil

Grease" Petrol Benzene Toluene l·\ethyl alchol A.cetone Trichloroethylene Sulphuric acid

(% aqueous) Hydrochloric acid

(5"/0 aqueous)

Test

o temperature C

20 50 20 50 50 20 20 20 20 20 20

20

20

7

Retention of

flexural strength %

99 98

100 100 ~~ ":-".1

99 92 94 98 96 97

100

99

1.2.1.1. MOLECULAR WEIGHT

Molecular weight (2) is measured from the usual relationship

[ '1 1 -a

= KM;;_

where [1 1 is the intrinsic viscosity (IV), K and a are constants

depending on the solvent, and Mii is the number average molecular weight.

The main effects of molecular weight on mechanical properties can

be summarised as follows:

For a fixed temperature at a given level of crystallinity, yield

stress and modulus show little dependence on molecular weight, but

they do show an increase with increasing crystallinity above a

certain molecular weight.

~ 2 Yield strainAwith crystallinity, but has no dependence on molecular

weight.

But, 3 An increase in intrinsLc viscosity (ie molecular chain length)

gives a correspondingly higher impact strength both in amorphous

and crystalline samples. The typical intrinsic viscosity for bottle

grade PET is 0,.74 dl/g. The principal cause of intrinsic viscosity

drop is hydrolytic degradation of the polyester chain. In the melt

state the attack of water on ester linkages is rapid and quantit-

ative. Hydrolytic degradation will be discussed later.

1.2.1.2. CONTAMINANTS

The final quality of a synthetic polymer is highly dependent on the

quality of its monomer.

8

For PET production ethylene glycol is easily purified and highly

purified terephthalic acid is now a commercially available product.

Another way of ensuring PET purity is to take care in processing. PET

produces acetaldehyde in very small amounts when manufactured or moulded

(3). The level of PET grades after manufacture is usually controlled to

<3 ppm. In bottle preforms the concentration is between 4 and 9 ppm

depending on the bottle type. The thermal decomposition to yield

acetaldehyde does not lead to any significant intrinsic viscosity loss.

It has been reported by Zimmerman (4) that when acetaldehyde is

retained in the polymer, a reaction is possible leading to a polyene and

water. The progress in manufacture in recent years to produce low acetal­

dehyde content PET means that initiation of degradation at the.chain ends

is unlikely.

Small quantities of impurities or additives can interact markedly

to change the thermal properties of PET. In preparation and during

processing PET is subjected to temperatures of 270-290o C. The rate of

degradation at this temperature depends on the amount of metal compounds

present. Derivatives of Ca, Mg, Pb, Co, Zn, Mn, Sb, Ti and Ge are all

used to an appreciable extent as transesterification and polycondensation

catalysts (4,5).

~Small quantities of transesterification catalysts can accelerate

the first step in thermal degradation of PET which is the random scission

of the ester bonds which leads to the formation of vinyl ester and

carboxyl end groups. This reaction leads to a decrease in solution

viscosity. It is necessary to produce pure PET to minimise the occurrence

of such reactions.

9

1.2.1.3. WATER

The principal cause of intrinsic viscosity drop is hydrolytic

degradation (3,5,6). When water is present in the PET melt the

polymerisation is driven backwards breaking up the large polymer chains

into smaller units. In the melt state the attack of water molecules on

ester linkages is rapid and quantitative. The degree by which any

starting intrins.ic viscosity will be decreased can be related to resin

I moisture content, Fig. 1. For example shown by the broken line is a

0.74 IV resin containing 20 ppm moisture; in processing it suffers a 2%

loss of IV to 0.725. This represents an acceptable loss. Physical

properties drop off significantly at MW level < z 15 000 (Mn) corresponding

to moisture levels above 0.02%.

MOISTURE

(%)

0.04

EFFECT OF MOISTURE LEVEL ON PHYSICAL PROPERTIES

PROPERTIES COMPARED WITH THOSE OBTAINED FOR DRY SPECIMENS

TENSILE STRENGTH

(%)

94

ELONGATION

(% )

90

UNNOTCHED IZOD-IMPACT

STRENGTH

(%)

64

It can be seen that the moisture content of 0.04% drastically affects

the impact performance of PET; drying is very important in the processing

of PET and this is discussed later.

Zimmerman and Kim (5) provide evidence that the hydrolytic degradation

of PET is autocatalytic, ie the carboxyl group concentration causes

acceleration.

10

REsr:; IV

0.8

- - ?- - - - - --O '7 .,

0.6

TIHE

0.0

2.5 -._--. 10 pno,-- H.p

5.0

j 7.5

10.0

~.~ I.OS::l OF IV

TIHE ?IG.1',IVloS3 of P~7' melt due to moisture ir: !"B.si:-~ (:;).

(r,fte!" ~ich"ul)

11

Thus hydrolysis of PET is dependent on the initial carboxyl group

content. Low carboxyl group content is crucial for thermal and

processing stability. It is also important for hydrolytic stability.

PET exposed to water for an extended period of time shows a noticeable

decrease in degree of polymerisation.

1.2.2. GLASS FIBRES

The basic concept of fibre reinforcement is the 'production of a two­

phase composite structure in which the matrix is used to transfer stress

by shear at the fibre-matrix interface, to the embedded, high stiffness

and strength fibres. Observations suggest that in terms of the commercial

exploitation of high modulus fibres in thermoplastics there is no advan­

tage in using any high performance fibre other than the cheapest eg glass.

The strength of glass fibres can be exploited by binding the fibres using

a resin or a plastic. Providing the length of the fibres is sufficient,

they should be constrained to take up the same deformation as the matrix

over the greater part of their length and thus effectively reinforce the

matrix. In addition, the presence of fibres often helps retard the prop­

Qgation of cracks and thus produce a material which is tough as well as

of high strength, but it does depend on the matrix.

Glass-reinforced plastics are composite materials (7), previously

they consisted of glass fibres in a plastic of the thermosetting type.

Now thermoplastic materials are used as the binder; the advantage of the

thermoplastic, its' ease of fabrication is retained. Most of the resins

used for the manufacture of glass-reinforced plastics are polyesters, also

used are polyamides and polypropylene.

12

In plastics the low modulus of the resin allows for the effective

transfer of stress to the high modulus fibres.

Generally (8), the addition of glass fibre leads to substantial

improvements in strength and rigidity accompanied by a good impact

strength, extremely high heat deflection temperature, excellent dimensional

stability, very low creep and superior long-term wear characteristics.

Glass filaments represent a low cost (9), readily available reinforce-

ment. Prices for glass reinforced materials typically range from £2-4/kg.

Glass fibres are less attractive than carbon fibres in applications that

require a high strength-to-weight ratio due to their high density (up to

3g/cm').

The most widely used type of glass for fibre manufacture is known

commercially as 'E' glass it is a soda-free calcia-alumina-borosilicate.

-1 This is drawn from the molten state at speed of up to 2 x 10'cm",.s to

diameters of between 5~m and 20~m. The tensile strength of the glass after

-2 -2 drawing averages 2.86 GNm ,with a Young's modulus of 70GNm ,a Poissons

ratio of 0.22 and a specific gravity of 2.55. The composition and

properties of some glasses are shown in Table 4.

The stiffness of short fibre reinforced thermoplastics depends on the

volume/weight fraction of fibres, fibre length, the stress transfer

efficency at the interface and the orientation. Similarly the strength

of a short fibre reinforced thermoplastic is dependent on fibre length,

weight/volume fraction of fibres, the interfacial shear strength and

fibre orientation.

13

TABLE 4: COMPOSITION AND PROPERTIES OF GLASS FILAMENTS (7.8)

Name

E

A-Soda-Lime

C

M

s

54.4

72.0

65.0

53.7

65.0

fused Silica 100

1> 2.54

2.50

2.54

2.89

2.48

2.2

14.4

0.6

4.0

25.0

E 72,4

69.0

112

85.6

73.1

CaO

17.5

10.0

14.0

12.9

v 0.2

0.16

Compositions are in weight per cent.

MgO

4.5 8.0

2.5

3.0 5.5

9.0

10.0

n a 1.548 4.9

1.512

1.541 4.0

1.635 3.2

1.523 1.6

1.458 0.31

BeO

8.0

T, 616

695

552

760

1070

14.2

8.0

ULTIMATE TENSILE STRENGTH

3400

3200

3100

3400

4800

p = density in 103kg m- 3j E = Young's modulus before compaction

in GN m-2 ; v = Poisson's ratio; n = refractive index at 550 nm;

a = linear expansion ccefficient at 25~C to 1000e in 10-6 per °C

T1 = temperature in °C at '.,;hich viscosity is 1014 . 5 poise

Contains also 3 per cent ce02

and 2 per cent Li2

0

Ultimate tensile strength in MN/m2

Strain point in K (softening Temperature of glass filament)

Source Kelly and Catherall

14

0.5

0.5

2.0

STRAIN POINT

889

800

825

800

1033

0.4

7.9 0.5

1.2.2.1. FIBRE CONTENT

The incorporation of fibr€$ is generally easy and fibre loadings

up to about 70% are realisable. When thermoplastics are reinforced

with glass fibre, the effective glass loading often represents a

compromise between the mechanical properties and the surface quality

required for the mouldings. Figs 2,3, and 4 show clearly that for

PBT (and similarly for PET) the modulus of elasticity, tensile strength

and notched impact strength rise with increasing glass content. It can

also be seen that the tensile strength and notched strength do not in-</<;

crease over 50% glass fibre, therefore higher loadings than this are of

no advantage. It should also be noted that a glass content of only 20%

by weight suffices to raise the deflection temperature to 200 0 C.

Widely available are glass-reinforced grades of PET with 10-50% by weight

of glass fibres. A linear dependence of strength on volume fraction of

fibres is expected from the equation

IT uc = (Tuf v (1 - ;~) +

, (l-V)crm for l>lc

lc = critical length.

·~f . ~ She(fi, of {ib~ L>:. = fcM.Ju>e. s~ 'd ~ile v . • ,f" ..... -f.t<u,"",,"

o-M : SNSS ~cI ':J /I..,. ""a-tN by Lavengood (10) and Blumentritt (11,12) This has been observed

for volume fractions covering the range 0-50%. At high concentrations,

mutual interaction between the fibres can result in loss of fibre strength

and excessive fibre breakage.

1.2.2.2. FIBRE LENGTH

E-glass average length = 200~

average diameter = 10~

15

of

:~su~e 2 Modulus of elasclci~y (tensile test) as a f<Jnc t io!: 0::'" ,.::1:.15S c on ten t (JIl·; 53457) for Crastine (?E~ reinforced with short ;lass fibre) at 23°C

12000

elastici tj~

o 10 20 50

Glass fibre content ,/? ~. 3 fl mm - ! l.gure

. 200. Tensile. st~~r.~t~ 38 a :unction of glass content (DIN 53455) for Crastineo (?3T reinforced with Ghort glass fibre) at 23 C

Tensile 150

st~enGth

\otcl1ed

i.mpact

;;tren~tn.

100

50

. _, 2 r~v / m

16

12

4

o

1

o

10 20 30 fibre conte!lt

Fi?;ure !.;.

4(: 50 (7~' by wt.)

Kotched impact strength as ~ func~.ion of glass content· for Crastine (PBT reinforced with short ~laRB ~ibre) at 23°C

10 2C 30 Glass fibre co:!te::t

50 ::-:. )

16

The glass fibres exert their effect by restraining the deformation

of the matrix (95). The theory is that the external loading applied

through the matrix is transferred to the fibres by shear at the inter­

face. Complex stress distributions in the fibre and matrix are the result.

For short fibres there is an increase in tensile stress from zero at the

end of the fibre to the value,~ f (max), which would be achieved if the

fibre were continous over the whole length of the matrix. To permit the

full load carrying potential of the fibres to be realised the area of the

interface must be sufficient, since the load is transmitted from the

matrix to the reinforcement via the interface. This means that for a

given diameter of fibre there is a length that must be exceeded if the

composite is to fail through tensile fracture of the fibre rather than.

shear failure at the interface. This fibre length is called the Critical

Fibre Length lc. Long fibres give rise to high stiffness and tensile

strength and also improved mechanical properties when exposed to elevated

temperatures, high continuous service temperatures, and high creep

rupture strengths under static loading.

The advantage of grades reinforced with long glass fibres can be

seen in the time-to-failure curves, Fig 5. These curves illustrate the

relationship between the applied load and the time to failure. They

show that Crastine grades SG623 and SG625, which are both reinforced

with long glass fibres, have about 30% higher creep strength in flexure

than comparable short-fibre grades. To achieve high values for stiffness

and strength (13) of a short fibre composite, it is necessary to use

long fibres, well bonded to the matrix. For maximum toughness, on the

other hand, it is desirable to have a weak interface or use fibres

having l(lc. Fibres should be of sufficient length compared to their

diameter to ensure an effective transfer of stress from the matrix to

17

?lexural

strength

Figure 5

Time-failure lines (flexural teGt) !or Crastine~SG 625 (1), SG 623 (2), SK'605 (3) and SE 603 (4) at 23°C.

,/ 2 N, !!lO

200E! ;;;;~~~~~Il 160 I J'32

120 I

80 I I' 1I

40 , I i 0~1 ________ ~----------+,--------~

10 lOO

"'. _~rn.e

1000

.& ~~ ~' ~ C.J:n - ~ew Ltd.

Sc.r - f'e.T ~orud ""iih. J.o~ ~ twttb SI( - fBT ~0rt.0..Q .wiiJ,. $ho.1; ~ Ji.b~

18

the fibres.

1.2.2.3. COUPLING

Coupled glass-fibre reinforced PET is an injection-moulding

material which is available in a number of different grades. The grades

can differ in the method chosen for chemical coupling, in the amount of

glass fibre present, and in the compounding method used. The distinction

betwen glass-fibre reinforced PET and the uncoupled materials is important

since major differences exist between their properties. It is important

to point out these composites are reinforced not filled, ie there is a

bond between the polyethylene terephthalate and the glass fibres which

enhances strength and stiffness properties. As a result the superior

properties of coupled grades make their use highly desirable in load­

bearing applications. Coupling agents are molecular bridges between

polymer. Titanium derived coupling agents (14) have a unique reactio~

with free protons that create polymer compatability, organic multi­

functionality and novel surface energy modifications. Titanate-derived

interfacial chemistry opens a new era in filled polymers, allowing un­

precedented incorporation of high filler Ifibre loads with conc.oMI!-attI:.· improve­

ments in polymer physical properties and rheology.

19

1.2.3. IMPACT MODIFIERS

There are a number of impact modifiers useful for polyesters

including polycarbonates, and polyester blends. Thermoplastic

moulding compositions usually have their impact strength improved by

incorporating an acrylic or methacrylic grafted polymer of conjugated

diene or an acrylate elastomer, alone or copolymerised with vinyl

aromatic compound. Especially preferred are the polymers available

from Rohm and Haas, for example, Paraloid KM653, Paraloid KN030 and

Paraloid KN611. ,-

~

It has been suggested (69) that the impact strength of thermo-

plastic polyesters can be improved by incorporating low density poly-

ethylene and glass fibres, particularly for a composition of PET, with

an aromatic polycarbonate and 5-50% by weight of glass fibres.

Linear low density polyethylene when added to an. aromatic PC results

in moulding compositions having improved weld line strength and heat

stability while retaining their good impact strengths. Further it has

been disclosed that compositions comprising. LLDPE and PET/PC blends have - ---------

improved com~~tability, weld line strength, flow properties and mould

releasability, including reduced plate out. I G~) ")

It has been discovered that thermoplastic polyesters show improved

impact strength when, preferably,5-15% by weight of LLDPE and 5-50% by

weight of glass fibres are incorporated. There appears to be a synergism

betwen the LLDPE and the glass fibres accounting for the improvement.

20

?

1.2.4. CRYSTALLISATION NUCLEANTS

The low crystallisation rate and slow nucleus formation of PET

constitute a serious handicap in injection moulding. Injection into

heatul .. moulds (150°C) to accelerate crystallisation and to shorten

cycle times yields finished articles of low crystallinity that are

difficult to remove from the mould and are too soft; if very long cycle

times are used the resultant articles are too brittle (due to large

spherulites). These problems can be avoided by the use of nucleating

agents.

A high degree of crystallinity and high nucleation rates can be

. achieved with high growth if heterogeneous nucleation is employed.

Nucleating agents have the overall effect of promoting rapid freezing,

giving a high degree of crystallisation, and reduce skin effects and

formation of voids which can occur in conjunction with large morphological

structures (70).

The primary nucleation of polymeric materials is of

considerable technological importance, it regulates the number and size of

spherulites which in their turn influence the impact resistance and

ultimate tensile strength and elongation at rupture of the material.

Primary nucleation depends on several factors, such as the melt treatment

prior to crystallisatio'n, the thermal pretreatment near the glass trans­

ition (Tg) and the presence of heterogeneities.

Table 5 shows the wide variety of nucleating agents for PET. In

practice, insoluble, inorganic nucleating agents such as metal OXides,

metal salts, and certain materials with particle sizes 3~m are preferred

21

Table 5: Nucleating agents for polyethylene terephthalate (71)

INERT, INSOLUBLE SUBSTANCES

Mineral fillers such as chalk, gypsum, clay, kaolin, mica, talc, silicates.

Pr~phyllite

Pigments such as cadmium red, cobalt yellow, chromium oxide

Metals: metal oxides such as titanium dioxide, magnesium oxide,

, a~timony trioxide; phosphates

Carbonates and sulphates, preferably of the alkaline-earth metals

Boron nitride

Sodium flVoride

Carbon black

ORGANIC COMPOUNDS

Salts of monocarboxylic or P01YCarbo~/ic acids

Montan wax and montanic ester salts

Diphenylamine, tetrachloroethane

Acetone, nitromethane, benzene, toluene

Halogenated Alkanes, such as tetrachloroethane

Benzophenone, tetralin

Aromatic alcohols and amines

Alkali aralylsulphates

[poxides

POLYMERS

Polyolefins: PE,PP, poly-4-methylpent-1-ene, poly-3-methylbut-1-ene

Copolymers of ethylene and unsaturated carboxylic esters

Ionic copolymers of ethylene and salts of unsaturated carboxylic acids

Copolymers of styrene derivatives and conjugated dienes

" (After Gachter and Muller)

22

(71,72) and used in concentrations of about 0.5%. They can be added

before, during, or after polycondensation, in the form of dry, fine

powders or in suspension. Tumble blending is another possibility.

Talc, kaolin, sili.ca and titanium dioxide have been used as

fillers (73), they act as effective nucleating agents for PET. The

overall rate of crystallisation depends on the volume concentration,

the size distribution, and the nucleating ability of the additives.

The occurrence of trans crystallinity is attributed to extensive

heterogeneous nucleation induced at the filler surface. From the

shape of the crystallisation isotherm, it can be concluded that the

crystallisation depends on the type (size) of the filler. The

crystallisation of PET as a function of the nucleating agent used, the

substantial reduction of the crystallisation time and the nucleating

efficiency can be seen in Fig 6. The nucleating efficiency decreases

in the following order talc>TiO,>SiO,: In the temperature range con­

sidered talc is the most efficient heterogeneous nucleating agent.

Nucleants can be selected from the group consisting of monomeric

esters of citric acid and epoxised esters of unsaturated aliphatic

carboxylic acids (74); these additives promote crystallisation rate

and improve surface appearance in polyethylene terephthalate.

Du Pont (75, 76) have reported that alkali metal salts of organic

acids and low molecular weight organic esters of aromatic carboxylic

acid improve surface appearance of PET. Oligomer·re.:. polyester

plasticizers (77) in polyethylene terephthalate compositions improve

crystallisation behaviour.

23

Figure 6 Nucleating ef:iciency of heterosen.~s nucleatin; agents (73).

Half-time of

crystallization

Crystallized

fraction

lOOt-

~

10

6. Unfilled· PET • PET

o PET .. PET - Kaolin

o PET - Talc

2~----~--~~--~------100 105 110 115

Temperature (oC)

-' :L~ --~-------

I

o.il[ I 0.6 .

0.4

0.2

o

6 Unfill~d-PEj

• PET- li02 o PEl - Si 02

:t PET - Kaolin

o PET - Talc

Time (min)

(Af~er Groeninkx et al)

24

1.2.5. POLYCARBONATE

Since its introduction in 1957 polycarbonate based on bisphenol A

has become firmly established as an engineering plastic. Polycarbonates,

in general, have excellent properties compared to many other engineering

plastics. These include; impact strength, creep resistance, broad

temperature range of use, dimensional stability, clarity and hardness,

rigidity and abrasion resistance. PC has superior ductility and impact

beaviour compared with all other amorphous plastics.

Gnauck (15) studied PC in aqueous solutions of salts and organic

compounds. PC proved to be stable to stress cracking at 20°C with

respect to all aqueous solutions ie the cracking induced flexural modulus

is above 30GN/m'. PC is largely stable to water, aqueous salt solutions

and aqueous solutions of formaldehyde, urea and resorcinol. PC is stable

to weakly alkaline solutions, but not to ammoniacal aqueous alkali sol­

utions which result in formation of stress cracks or attack the surface

but result in no marked change in mechanical properties. PC is relatively

stable to aqueous solutions of mineral acids and carboxylic aCids, but

concentrated acids cause considerable deterioration of mechanical properties.

Plastics exposed to organic media may fail at stresses much less than

their yield stresses (16). Broadly termed environmental stress cracking,

this phenomenon manifests as a reduced service life time for many

plastics and as the rapid crazing and cracking of strained glassy plastics.

The use of PC in fields such as the automotive industry is hampered

by their poor resistance to petrol and other fuels (1, 17). Critical

s trains for PC range from· 0.33 to 0.78 per cen t.

25

In general the lowest critical strains were observed using petrols

which are severe cracking agents, critical strains above 0.4 per cent

were observed for petrols which caused crazing of PC.

A general correlation was observed between the aromatic content

of the fuel:- and critical strain. Critical strains for PC decrease

as the aromatic content of the petrol increases. For PC critical strain

initially decreases with the increasing toluene (or aromatic) content up

to thirty per cent and levels off at higher concentrations. PC exhibits

higher critical strains when exposed to aliphatic components than when

exposed to the aromatic components of petrol. Because of the low values

of critical strain for the pure aromatics, the effects of aromatic

structure on the critical strains for PC were measured by a continuous

drip procedure and were relatively constant over the entire range of

solvents. Splashes of mixtures containing high molecular weight components,

such as 1, 2, 4 trimethyl benzene and n-butyl benzene, resulted in severe

cracking, while continuous exposure in the same mixtures caused only

crazing;:_ The increase in severity of cracking and the reduced critical

strain were suspected to be due to the preferential. evaporation of the

non-aromatic component from the mixture. Critical strain measurements

give a reproducible estimate of the type and severity of the type of

petrol-induced stress cracking of complex moulded parts.

Another problem with PC is its hot water ageing behaviour (18).

The solubility of water in PC increases from about 0.3 per cent at

room temperature to 0.6 per cent at 100°C. However, these relatively

26

low concentrations of water have significant and even dramatic effect

on the mechanical performance. Hot water causes gradual chemical

degradation of the polymer; the long term use of PC in water having a

temperature above approximately 60°C is not recommended.

Table 6 shows the effect of some solvents on PC. When the solubility

parameter of the liquid falls into the inert regions where there is no

effect on the polymer as far as crazing and fracture are concerned. When

the solubility (94) parameter is in the range of ·9 to 10.7 approximately,

it is a cracking liquid for PC causing dissolution and fracture.

Blending is thoughtto improve the chemical resistance of PC (1).

1.2.5.1. MOLECULAR WEIGHT

The mechanical properties of aromatic polycarbonates depend on the

molecular weight. The molecular weights (Mn) of polycarbonates generally , -

fall in the range 35,000 to 40,000 and in this range the polymer will

have moderate mechanical properties. Above this range the values for

impact and tensile behaviour significantly i~crease. The rise is not

limitless, however, above a certain value the properties level off.

27

TABLE 6: Solubility parameters of liquids (94) " ~

LIQUID Solubility Remarks

parameter . - "): .J./

(call cm ~)L

Ethanol 12.92 12.7 Crazing Agent

Butanol 11 .30 11 .4 Crazing Agent

Cyclohexanol 10.95 11 .4 Crazing Agent

Methyl carbitol 10.72 Crazing Agent, polymer attacked

Hexane 7.24 7.4 Crazing Agent

Cyclohexane 8.18 8.2 Crazing ·Agen t

Butyl acetate 8.46 8.5 Cracking Agent, polymer severely attacked

Dibutyl phthalate 9.3 Cracking Agen t, polymer sevetply attacked

Water 23.5 23.4 Polymer not affected

(After Miltz)

28

1.3. PROCESSING

1.3.1. DRYING

Both 'polycarbonate and polyethylene terephthalate have absorbed

water which must be removed before melt processing to prevent hydrolysis.

Physical properties, particularly impact strength, deteriorate significantly

at moisture levels above 0.02%: properly drying the polymers is the first

crucial step towards obtaining high quality parts. Parts moulded from

wet PET do not exhibit surface defects, so that parts could be moulded with

excellent surface appearance and yet have poor end-use performance.

Drying the material is relatively simple' (6). Virgin PET and regrind

must both be dried to less than 0.02% moisture and kept at that level for

processing. PET resin picks up moisture very rapidly; for this reason the

use of remote tray-oven dryers with manual transfer to the hopper is not

ideal. In the case of regrind it is important to keep the grinder blades

sharp (23), and to adjust the clearance between screen and blades in order

to minimise fines. Fines have more surface area than the pellets and

pick up moisture faster.

PET and PC are hygroscopic - ie moisture is not only collected on

the surface but it is absorbed inside the pellets. For hygroscopic resins

hot, dry air is required for pulling moisture out of pellets/regrinds.

This requires a dehumidfying (dessicant) type dryer.

Recommended conditions are:

1. Air flow 0.05 to 0·010 !1'I~/min for each k.g per hour of resin.

29

2. Air temperature is important, it should be measured at the entry

point to the hopper. Temperature should be 135°C.

3. Drying times depend on the moisture content, 2 hours for a resin

with a 0.04% moisture content. For very wet resin drying time

should be extended to 4 hours; however, prolonged drying at 135°C

is not recommended. If the resin is being dried overnight (longer

than 4 hours) drying temperature should be reduced to 107°C.

4. Dew pOint of the air entering the hopper must be OOG or lower

throughout the drying cycle in order to dry the resin adequately.

1.3.2. PRECOMPOUNDING

Nassar et al (24) carried out blend preparation in the following way.

Pellets were dried and the hot pellets were transferred to the mixing

bowl which had been preheated to 260°C. This charge. when melted/completely

filled the mixing chamber. The mixing blades were set at the speed of

2-4r.p.m. during polymer addition. After all the pellets had been added,

the lid was closed to minimise oxygen absorption by the polymer, the speed

of the mixing blades was raised to 90 r.p.m., and the heaters adjusted to

obtain a final blending temperature of 290-300°C to ensure melting of the

crystallin~ PET. The high speed was employed to reduce the mixing time

needed to approximat~\y 8-10 mins. Compounding is referred to more gen-

erally in 1.3.4.

1.3.3. INJECTION MOULDING

Composite injection-moulding compounds consist of short fibres

dispersed in a thermoplastic matrix. Injection-moulding compounds have an t~fO~

advantage over short-fibre sheet moulding ~and continuous fibre systems

30

because of the possibility of moulding complex shapes (19). The j"

principal disadvantages are:the relative~~.soft matrix)fibres break

under high shear and a lack of predictability of the ultimate properties.

Since the properties of a short fibre reinforced thermo plastic are

very depend.nt on fibre length and orientation, it is important that both ,

of these parameters can be controlled in the final moulding, by an approp-

riate choice of processing conditions (Table 7).

TABLE 7

Normal conditions for Rynite+530 and 545

Melt Temperature 295 + 5c C

Preferred Mould Temperature 90 - 110°C

Injection Pressure 10. - 90 MPa

Injection Speed Moderate to fast

Screw Speed as low as practicable

Back pressure none or as low as practicable

-RYNITE 530, a general purpose grade containing 30% by weight of glass

reinforcement, offering outstanding strength, stiffness, toughness and

appearance for moulded parts

-RYNITE 545 - as above, containing 45% by weight of glass reinforcement.

Although the use of slow screw speeds, slow injection rates, low

back pressure, wide sprues, runners and gates, and large radii of curv-

ature minimises fibre breakage during moulding, such conditions are not

31

found often in practice. Furthermore, the necessity of incorporating

reground material into the feed-stock also ensures short fibre lengths

in the final part, lengths not greatly in excess of the critical length

required for effective transfer from polymer matrix to reinforcing fibre.

The process conditions for moulding glass-fibre reinforced PET

were determined by Haworth (20) using a Bipel machine (Table 8).

Table 8

Melt temperature (settings) 260/268/274°C (barrel)

(nozzle)

Mould temperature 140-150 o C

Injection pressure (melt, maxI 2670 psi (18.4 MNm- 2 ) (Izod)

Injection pressure (melt, maxI 5723 psi (39.5 MNm- 2) (Tensile)

Injection/Hold-on time

Cooling time

Screw rotation speed

1.3.3.1. MOULD TEMPERATURE

20s

50s

65 r.p.m. (1.08rp5)

The melt temperature used for reinforced thermoplastics is usually

at the upper end of the range recommended for the unfilled counterpart.

This is chosen to reduce the viscosity of the melt and to assist in

preventing premature solidification in the cavity.

The hottest part of the mould is immediately opposite the feed point

so cooling should be concentrated there.

32

Tubes, pipes and channels should be uniform in cross-section, to avoid

variable flow rate in the coolant and the creation of hot and cold spots

(21). The flow of the coolant should be from bottom to top so as to

keep the system full. The high heat transfer, metals of high thermal

k-' conductivity eg beryllium and copper (conductivity 290 W. ~-, ) are

commonly used for selected parts of the mould.

It should be emphasised that running the mould at the lowest possible

temperature is not necessarily conducive to obtaining the highest quality

mouldings. Mould surface temperatures between 85 and 120'C are suggested

for optimum dimensional stability, surface appearance of moulded parts and

cycle time. The preferred range is from 90 to 110°C. When low mould

temperatures between 60 and 85°C are used, the initial warpage and shrink-

age will be lower, but the surface appearance will be poorer and the dim-

ensional change of the part will be greater when the part is heated above

85°C. If a minimum as-moulded warpage is the only requirement, resins

should be processed with mould surface .. temperatures of less than 60°C.

1.3.3.2. GATE DESIGN

By properly selecting the flow characteristics of the moulding compound,

the moulding conditions, and the type of gating, the orientation dist-

ribution in the moulding can be controlled. Design is especially important

for reinforced thermoplastics, there is the obvious need to locate the

gate(s) at realistic positions in order to ensure development of an approp-

riate fibre orientation distribution in the final component. The highest

shear rate is in the gates, therefore there is much fibre alignment.

33

The gate dimensions must also be larger than would normally be

appropriate for unfilled thermoplastics. This is due to the likelihood

of jetting occurring during the filling of the cavity and the need to

ensure that prolonged hold times can be used, without gate freezing.

For this section moulding, it is necessary to use a sprue gate or pul­

sating injection otherwise unacceptable voiding will occur in the core

of the moulding.

Orientation produced by fan gates (22) depends strongly upon the

fill rate and may be considerably more transverse than that produced

by edge gates. Among edge gates, the largest produce a somewhat larger

component of orientation in the longitudinal direction. This orientation

is about equal to the highest level obtained with the fan gates. Tunnel

gates can be used, provided the gate diameter is greater than 0.5mm. In

three plate moulds, the gate diameter should be about 45 to 75% of part

thickness. For rectangular gates, the gate thickness should also be

about 45 to 75% of the thickness and the gate width should be 1.5 to 2

times the gate thickness, depending on part volume and surface appearance.

For both round and rectangular gates, the gate land should be short:

between 0.5 and 1 mm.

1.3.3.3. INJECTION SPEED

High injection speed should be used in order to achieve a good

surface finish to prevent premature solidification of the melt, either

in the cavity or at the gate. The screw speed and back pressure must be

kept to a minimum, since alth.o"3h-:· a homogeneous melt is required, fibre

breakage may become excessive.

34

1.3.4. COMPOUNDING - ADDITION OF FIBRES

The selection of the appropriate compounding technology is very

dependent on the requirements of fibre length, volume fraction and

degree of dispersion of the fibres throughout the matrix. One of the

most common methods of compounding involves the use of a single or

twin screw extruder to mix chopped fibres and matrix and produce an

extrudate, which can be pelletized to give roughly spherical granules.

Twin-screw extruders provide considerable advantages for processing

highly viscous materials (25). A specific advantage and distinctive

difference with respect to single screw extruders is the self wiping

design and consequent elimination of stagnant zones. Twin-screw

extruders can be co-rotating or counter-rotating.

To produce high grade reinforced compounds processing is usually

by means of a twin-screw extruder (26, 27) which can perform other­

wise complicated compounding relatively easily. The twin-screw can

prepare fibre-reinforced polymers with a minimum of breakage/damage

and also blend polymers of differing chemical composition and physical

characteristics.

Compounding extruders must incorporate special design features

to cope with the abrasive nature of the glass fibres, the high powers

required for compounding and possibly novel methods for feeding the

polymer and fibre. The machine wear, especially on the screw and

barrel, can be quite unacceptable unless special nitrided steels or

hard alloy coatings are used.

The high shear stresses generated in extrusion compounding

equipment for effective dispersion often result in severe damage to

35

brittle additives. In the case of fibrous inclusions such as

glass,fibre, a significant reduction in aspect ratio may result,

which may greatly reduce reinforcing potential in a thermoplastics

matrix.

Work has been carried out at Brunei University which demonstrates

that mean fibre length is reduced most dramatically in the early

stages of the twin-screw compounding operation where the polymer is

in a solid or semi-molten state, see Fig 7. Feeding the fibres in at

the middle of the barrel provides a means of reducing fibre breakage;

the fibres are added after the polymer has become molten.

The twin-screw action readily takes up the fibre - a feature

that is not normally exhibited by a single screw machine, unless

special consideration is given to the design of the feedport. The

addition of fibres to a pre-melted,polymer has the advantage that

less fibre breakage occurs together with an improvement in dispersion.

There is the additional advantage that much less wear takes place in

the extruder.

1.3.5. PROCESSING EFFECTS ON PROPERTIES

1.3.5.1. CRYSTALLISATION

The rate of crystallisation in PET is of basic importance.

Since the polymer crystallisation is usually accompanied by evolution

of latent heat, differential scanning calorimetry (DSC) can be

36

Fibre length in mm

5

4

3

2

1

o

FEEDING

DiO:COHPR;;;S.sION

I-!ETERIKG

L DISTXNC3 ALONG SCR3WS

?ibure 7 The c~ar.~e in ~ean fibre length during twin-scre~ extrusion co~pounding. Results for glass fib~es added to poly prop­ylene ~e~onstrate the" considerable da~ace that can occ~r to the fibres in the early staGes of co~poundinG. (27;

37

utilized for following the course of crystallisation. Quantitative

analysis of crystallisation kinetics can be performed using Avrami

parameters (28).

The heat evolved during crystallisation yields exothermic peaks

in DSC. If DSC is operated under the isothermal condition, the

heat of crystallisation is obtainable by measuring the area under

the thermogram peak. If t max denotes the time to attain maximum

rate of crystallisation then it can be determined from the Avrami

equation.

X(t) = 1-exp [_ktn]

where k and n are constants.

From this equation can be derived (28) the relationship between- the

degree of crystallisation and Avrami exponent, n. It is also derived

that the temperature dependence has the following relation:

ln t max = a + b T'6T

where a and b are constants, and 6T = Tmo-T, where Tmo is the

equilibrium melting point of the polymer. This indicates the temp-

erature dependence of crystallisation rate. Generally it can be seen

in Fig 8 that a longer t is obtained at higher molecular weight max

of PET. This fact indicated that higher molecular weight causes ~

slowing down in the rate of crystallisation. A minimum value of t max

exists at around 175°C. Lin (28) also found evidence of the existence

of secondary crystallisation of PET.

38

4

3

2

1

Fig. 8

~jn = 36,000

200

(0, ... " , ~,

t vs TEMPERATURE (28) ma:x

(After LIN)

39

= 2~,000

2:::0

The DSC was also used by RoberlS (29) to follow the morpho-

logical changes whi~occur when partially crystalline PET is annealed.

Annealing would increase the perfection and size of these crystallites

by the formation of thin lamellae and bring aboub further crystal-

lisation of the amorphous polymer. This occurs by the initial formation

of crystallites having even larger surface free energies and hence

lower melting points than the intial imperfect crystallites. The

initial crystallisation of the PET brought about by cooling rapidly

from the melt produces imperfect spherulites and the overall morphology

is of the 'fringed-micelle' type. The individual crystallites are

small, have large surface free energies and low melting points.

When PET was examined using low angle X-ray measurements three

spacings were apparent; two at 12 ·13 "'" and 15, -16 nu being common

to all crystalline samples and an additional spaCing of much lower in-

tensity which varied between 20 and 50- M'·

()./ld T.vi SiegmanA(30) showed that the structure of PET could be altered

and more order gained at temperatures below Tg. It was shown that

annealing below Tg results in the formation at high initial rate of

an endotherm in the DSC thermograms, which continues to increase in

size for long periods of time. It also results in an increase in Tg

• + and Tc and a decrease in Tm. Annealing below Tg results in a material

which, upon annealing at elevated temperatures, crystallizes to a

higher degree and upon drawing gains more order but less orientation.

"c~S+alliscJicn !fr,.f~rrduN.

+ MeJ hllj ~ pero:fUJ!

40

As a result, the mechanical properties are less favourable.

PET is very sensitive to thermal history. Structural changes

which affect the polymer behaviour. and hence its properties,are

introduced by thermal treatments at temperatures below and above Tg.

1.3.5.2. FIBRE BREAKAGE/ORIENTATION

Compounding can result in a major reduction in fibre length.

The twin-screw compound extruder produces high grade fibre rein­

forced compounds with a m.j'n:imum fibre damage (31) to the reinforcing

fibres, (see Fig 9).

The high shear stresses generated in extrusion compounding

equipment for effective dispersion of fibres often result in severe

damage - a significant reduction in aspect ratio may result, which

can greatly reduce reinforcing potential in a short fibre reinforced

thermoplastic matrix.

Measurement of fibre length distribution in a polymer compound

can conveniently be made by ashing the compound. Figure 10 illustrates

a typical glass fibre length distribution curve obtained by semi

automatic counting of magnified fibre images, using a graphic table

connected to an image analyzer.

During the high-shear injection moulding process, considerable

fibre breakage can take place. This may result in fibres in the

41

:: l; .... ::re 9

!~ini~uc ?ibre ~reakaGe (31)

". ·'I:n.e i'~PC/~.; s:{ster;: p:-oduces high grade fi~re reinforced. compou~.c.s Hith a mini!num of breakarze/da:lsge to the reinforcin,5 fibres •

. RICHT·; HFC/V product; LEFT··.! p:oo:iuet f!"'.):;: E' .:::inf-le ser'?':: ext!,,1jc.e:- .-

42

50 ,.--

I-r-

40 -

f-

30

?::-ecuenc:.~ - -

20 I-

r-

-

10 . , r-

r-

r-t-

o I be[ -rh c 0.5 i.C

c;:dure. (27).

43

moulded component having length less than lc (typically <!! 200 IJ.m for

fibre matrix combinations) ie too small to ensure good stiffness and

strength of the composite.

Some modest degree of control is possible by manipulation of the

moulding conditions. The temperature at the feed zone of the screw

has a relatively large effect on fibre length. As the rear zone

temperature is increased there is an associated increase in fibre

length. The greatest increase in fibre length occurs when this temp­

erature approaches the melting pOint. This suggests that the outer

layer of the granules will',: soften and reduce the tendency for fibres

to break, as a result of intergranular interaction.

One major source of fibre breakage occurs during screw-back. This

part of the cycle is concerned with the production of a fully homogenized

melt and so normally, for an unfilled thermoplastic, this is aided by

employing a high screw speed and back pressure. These are not appropriate

conditions to use when fibres are present since excessive breakage will

occur. In view of these effects, it is normal practice to employ zero

or very low back pressure and the slowest screw speed consistent with the

cycle time.

Short-fibre reinforced thermoplastics have become increasingly

popular in many engineering applications since they can be processed by

conventional thermoplastics methods. However, the incorporation of fibres

such as glass, although significantly improving the load capabilities

of the materials, introduces, during processing, a complex inhomogeneous

fibre orientation distribution, (FaD), (32).

44

The orientation of the short fibres is determined by the flow

characteristics of the melt (33), which in turn depend on the mould

geometry, the wall thickness of the final parts, the characteristics

of the moulding operation, and the length and fraction of fibres in

the composite.

A knowledge of FOD is essential in seeking to predict the

deformation behaviour of short fibre reinfor,~dthermoplastics.

Examination of individual fibre orientation can be performed using

microtomed sections or surface prepared by metallographic polishing

techniques. These techniques are widely used since they are not

restricted to translucent mouldings, it causes little damage to fibres

and little disruption of their orientation. The process of examining

fibre orientation was made easier by the use of the technique of

contact micro-radiography (34). CMR is less time-consuming since the

specimen preparation by diamond saw is straight forward and semi­

automatic. Although a significant volume of the sample is imaged, all

fibres lying through the thickness of the slice are projected in focus,

onto the plane of the photographic film. A similar (improved) tech­

nique was developed by Hemsley (96).

Mechanical properties such as stiffness and strength can vary

considerably with changes in FOD, it is evident that as mouldings

become more complex it becomes increasingly difficult to predict the

FOD and hence the mechanical properties in any part of the component.

The moulding of thermoplastics involves non-isothermal flow ie

• the injection of a melt into a relatively cold cavity.

45

A skin layer of solidified material will form at the walls of the

cavity. The thickness of this layer will be affected by three major

variables, apart from the flow geometry:-

i. Melt temperature

ii. Mould temperature

iii. Injection Speed.

If, for example, there is a reduction in mould temperature it

will cause an increase in skin thickness, with the particular fibre

orientation in that layer. In general, sections taken through the

thickness of a moulded part reveal layers of differing fibre orientation.

In the simplest cases, a three-layer structure is found, with the fibres

in the core predominantly in a direction transverse to the main flow.

The fibre-orientation pattern in the two outer layers depends on the

mould geometry and the base polymer used. In a moulded tensile bar

of glass-reinforced PET over a large part of the section the fibres

tend to lie along the length of the bar. There is, however a prominent

central region of different less-aligned orientation; and a region of

lower alignment, which is more extensive towards the end of the bar.

1.3.5.3. BLEND CHARACTERISTICS

It was suggested by N~ et al (24) that processed blends of

PC and PET rich in PC exhibited two 'glass transitions while only

one Tg could be found from compOSitions rich in PET using either

differential thermal analysis or dynamic mechanical testing. These

results indicate miscibility over part of the composition range for

PC - PET.

46

They concluded that blends containing more than 70% PET by weight

form a single amorphous phase, whereas at lower PET levels two

amorphous phases exist. They do not offer any explanation why this

should occur.

In their second paper (35) the authors supported their earlier

conclusions, but emphasised that observations on crystallisation be­

haviour show very little evidence of interchange reactions between

PC and PET.

Birley and Chen (36), found·.that during extrusion the mel t flow

index (MFI) of pure PET increased indicating that thermal degradation

had taken place. The melt flow index of PC increased only marginally

indicating that its thermal stability was reasonable. Of particular

interest were the melt flow properties of different blends part­

icularly the 80/20; PETIPC blend which showed a compartively low MFI.

It was suggested this might be because in this blend some reaction

too~ place during melt blending and MFI measurement.

It has proved difficult to find the second transition temper­

ature of 80/20, PETIPC blend at about 140°C from DSC trace. A DMTA

spectrum of the blend however clearly shows that this blend has two

transition temperatures at 85°C and 140°C indicating that PC and PET

are not miscible in their amorphous phases.

The blends do not suffer from shrinkage or become brittle during

annealing at 125°C for 18 hours, and remain ductile, although the

80/20, PETIPC blend increased considerably in yield stress. The

improvement in the blend compared with pure PET is probably associated

47

with the PC in the blend retaining its softening characteristics

(ie Tg 140°C) and thereby stabilising the dimensions and modifying

crystallisation.

1.3.5.4. RECYCLING

The method developed for the recycling of PET seems suited to

producing degradation of the PET.

The major task is cleaning the scrap which requires large quantities

of water. The shredding and granulating of the bottles has the unfortunate

effect of increasing the surface area for maximum moisture absorption

when the granulated scrap is placed in the hydrocyclone. After this

separation stage the PET is saturated with moisture.

After the washing and separation stage, the PET is dried, melted

and extruded. The importance of this drying stage cannot be over­

emphasized, at 275°C a small amount of water can produce an enormous

amount of degradation. The PET must be dried to <0.02% moisture and

kept at that level for processing.

The drying method is also extremely important. Srter the washing

stage the granules will undoubtedly have moisture on their surfaces.

If the granules were then heated in an oven degradation would occur.

It is therefore necessary to remove the surface moisture under vacuum

prior to placing the granules in a drying oven.

The granules should be allowed to cool to room temperature in

48

the oven before transferring to the nitrogen filled hopper of the

extruder. The granules at room temperature take up a lot less

moisture than hot ones, and dry nitrogen in the hopper will force

out any air containing moisture.

The PET is extruded into a water bath and is then pelletized.

It is then stated (la, lb) that the pellets are crystallized, ie

heated in an oven until maximum crystallisation occurs. There is

no mention of a drying stage between pelletization and crystallisation;

this is a serious omission as an air knife cannot remove all the

moisture from the strand line, resulting in damp pellets. If these

pellets are then heated as implied a great deal of degradation will

occur. Care must be taken throughout the process if PET with properties

as close as possible to those of virgin material is required.

49

1.4. GLASS FIBRE REINFORCEMENT

1.4.1. POLYETHYLENE TEREPHTHALATE REINFORCED WITH GLASS FIBRES

The physical properties required for plastics used in the manu-

facture of functional components subjectto high loading can be attained

through modifying the base materials (37), in accordance with the

particular application. Fibre composites are formed by embedding the

fibre reinforcing component in a homogeneous material (the matrix) and

are distinguished by their anisotropic properties. Polymers are generally

reinforced with glass fibres (7).

In glass-reinforced plastics the composite material is put together

in order to exploit a property of the strong phase and a plastic is used

as a suitable binder. The tensile strength of glass after drawing is

-2 2.86 GNm . Very high ultimate tensile strengths can be achieved, for

_2" example, with E-glass (8) fibres when their surface is perfect: 2.50 GNm

is easily obtainable, with the modulus falling in the range 50-110 GNm _2

In some instances mor.e complete utilization of the properties of

individual materials can be made by combining them to form composite

materials (38, 39), so thermoplastics are increasingly reinforced etc.

to produce new materials that usually have enhanced mechanical properties,

compared to the thermoplastics matrix.

Apart from other advantages these randomly arranged fibre composites

are distinguished by their relatively good resistance to corrosion. This

is ultimately reflected in maintenance and servicing costs. There are

problems, however, (40), the design of many GRP structures is limited by

50

the loss of integrity which occurs at the onset of microcracking

at low composite strains, occurring particularly in those areas

where principal fibre direction is perpendicular to the direction

of an applied stress. Any flaws or microcracks expose fibres to

corrosive attack (41).

The majority of engineering structures are subjected to repeated

loading. Glass fibre reinforced plastics are becoming extensively used

(42) as structural materials, for machine parts and other material parts.

Polyethylene terephthalate glass fibre composites are distinguished

by their exceptional rigidity accompanied by good impact strength, high

heat deflection temperat"ure, dimensional stability, low creep and good

long term wear characteristics. In addition the composites possess good

electrical properties and excellent chemical resistance.

The compOSites can be processed in standard injection moulding

machines with attention to drying. These qualities give the composite

possibilities of numerous applications in many branches of industry eg

in electrical engineering, electroniCS, precision engineering, machine

construction, ship building and due to their ease of processing of

growing importance in the automotive.industries. (But, these are

presently limited to short fibres only, as breakage takes place during

processing. )

1.4.1.1. STRENGTH

Strength is conventionally the stress level at which failure

occurs (44).

51

Failure can be gradual or rapid and mayor may not be catastrophic in

nature.

Composite strength is given by (13):-

0" uc = O"uf v ,

+ (1 - v)crm for l>lc

where

0- uc = failure stress of composite

cruf = tensile strength of the fibre ,

(Ym = stress carried by the matrix at the fibre

failure strain

lc = CRITICAL FIBRE LENGTH

v = volume fraction of fibres

From the equation above a linear dependence of strength on volume

fraction is expected and has been observed (10, 11, 12) for volume

fractions 0 to 50% (see Fig 11).

The strength of composites reinforced with randomly orientated

fibres increases linearly with volume fraction up to 0.5. For all

loading levels, the strength of a random composite is somewhat greater

than that of the equivalent composite in which the fibres have been

aligned by flow processing. The random composites are superior to the

aligned system for all strength limited applications (nJ.JlcloM. LoCA.d.s).

52

Tensile

3trenGt:'·~

(1(31) K=10'lbs

SOt

I i

40~

! , , , ! I !

~L I i I ! i !

l0r-

.

..... : .' .

•• <

:::--, , \

10''---;;;--~';:----==--_".--_~_--,--10 20 30 40 60

VOL "10

?·2s/: r.:n-::!"'ice3 ;;nc: vc::,~rirr~ v:i2..ume . of ~las= fib~e :,ei~fnrce~e~t (10).

(~~ter ~~vengood)

53

There is a deviation from the linear dependence of strength

on volume fraction at low and high fibre concentrations. At low

fibre concentration matrix embrittlement is thought to occur promoted

by stress concentration at the fibre ends. At high concentrations

mutual interaction between the fibres can result in loss of fibre

strength and excessive fibre breakage.

From studies reported it would appear that the relationship

between strength and fibre concentration is more complex. It has

been customary to interpret the observed longitudinal strength of a

composite using one of the models describing composite failure. This

is because in most short fibre reinforced thermoplastics there"will be

a distribution of fibre lengths; hence it is necessary to sum the

contributions to the composite strength arising from fibres of sub-

critical and super-critical length:-

er uc .{I 1

Tu li Vi d

+ + (1 - V) a-'m (13)

The strength O-uc is related to the fibre length distribution and

interfacial shear strength Tu. If fibre length distribution and Tu are

known 0- uc can be calculated.

It has been found that calculated strengths considerably exceed the

experimental values. Most of this discrepancy could be accounted for

by internal stresses at the interface during the moulding and subsequent

54

cooling of specimens. The problem with reinforced plastics is

that they are inhomogeneous, anisotropic and rarely behave in a

linear elastic fashion (45). Crack paths are highly complex, and

the crack itself is not the only manifestation of structural damage.

Fibres break, the resin cracks, adhesion between the fibres and

resin may be destroyed and all of these processes will help to degrade

the mechanical properties of the composites.

In injection mouldeq thermoplastics the mode of cracking can be

complex, for example, if the predominant fibre orientation varies

through the thickness. However in many cases a relatively simple

dominant crack can be identified and its propagation rate measured.

If the fibres are sufficiently short, the crack may avoid the fibres

by local shifts in direction and by fibre debonding. In thin sections

of an injection moulded part, flow conditions can produce highly

anisotropic properties (46).

The fracture properties of pclyethylene terephthala te reinforced

with glass fibres were investigated by Friedrich (47). In the case

where-the fibres are perpendicular to the crack, the crack grows in a

fibre avoidance mode by-passing regions of agglomeration of locally I

aligned fibres. This leads to a zig-zag appearance of the crack path.

The L~cracks, (parallel to the fibre orientation), run straighter

and perpendicular to the applied load, utilizing the parallel orientated

interfaces between fibres and matrix. Both of these cases have been

illustrated in Figs 12a and 12b.

55

Fig~" 12 a Fatigue crack "iri" I Rynite 545-1-1 (45 w% fibresh

Fibres"perpendicu18.r to the crack direction eT-crack);

Fig."12b the crack (47)

Fibres parallel to direction (L-crack)

(After Friedrich)

56

In both cases the local crack tip advance occurs after wide­

spread formation of crazes, voids and microcracks ahead of the main

crack, mostly near fibre ends and along the fibre-matrix interfaces.

Close to the main crack tip these locally damaged areas coalesce be­

fore they tear apart to provide further crack growth.

Lhymn and Schultz (19) followed up Friedrich's work on failure

using fractography and linear elastic fracture mechanics to character­

ise failure.

When the fibres in a specimen are oriented along the direction of

flow, ie parallel to the surface (L-specimen) and a crack is growing

across the specimen, under an initially applied load the crack grows

to a certain length and then stops. When the load is increased frac­

tionally the crack propagates suddenly. Fig. 13 shows the step-wise

crack growth, the crack path has an irregular zig-zag shape but prop­

agates along the centre line on average (Fig. 12a).

Cracks start by the development of micrcracks at the fibre ends.

The bonding of the matrix at the interface results in the formation of

microcracks, which finally coalesce to form a continuous crack by

crazing of the matrix. Interfacial debonding follows the fibre end

microcracking. A large amount of fibre fracture is observed, induced

by the combined action of tensile and shear stresses.

Generally fibre fracture is completed long before the main crack

tip arrives. When the crack tip meets a fibre which is not fractured

?irru~e 13 Step-wise propa~atio~ of ~-craclts (19)

®

58

previously, interface debonding and fibre pull-out is the mechanism

of further crack development.

When the fibres in a specimen are oriented transverse to the flow

direction (T- specimen) and a crack is growing across the specimen;

parallel to the fibres (L-crack) the crack no longer has to grow by

fibre avoidance. The crack mechanism is mainly interface debonding plus

matrix cracking, a critical process being the crazing of the matrix to

join microcracks. Interface failure occurs far ahead of the main crack,

joining of these microcracks being the only remaining obstacle (Fig. 14).

This method of crack growth results in smoother fracture surfaces (Fig. 12b).

This continuously cracked plane occurs as a result of the matrix crazes

coalescing to form a crack which later tears apart.

The fibre alignment results in highly anisotrpic properties, the

two extremes are:

a) with the initial crack growth normal (90°) to the dominant

fibre orientation at the surface (T-cracks)

b) with the crack direction more nearly parallel (0°) to the

fibres (L-cracks)

The crack path at an angle between these can be predicted from the two

extremes. Table 9 shows tensile strengths for commercially available

PET/glass fibre composites; it can be assumed they use the most

favourable crack direction.

Lhymn and Schultz (19) performed tensile tests on PET reinforced

with 45% by weight chopped E-glass fibres. Pure PET was also tested

for comparison, showing an upper yield point and a large degree of

plastic flow before failure.

~i;ure 14 ?ar-field effect of a ~-speci~en (L-crack) (19).

2 ~icrocracks for~ at !ibre ends anc partial debondins by normal stress occurs

-:r:;::.;;;:::r: !·lain crack reac!1.es the fi~st ® "--:jiL,'2'.;:'.':;:"";: ... ".J'I: i' ·····,···'·,,1 -,-".--=cr fibre and debonding starts.

h:-:' .... -.... [! art i~ 1 ae bonding gro\rls in

QV ~-. --~I~'-"~:'-'~I~!:··':·:····~·'=·I[

(After Lhy~n t 5chultz)

! ',. :··:,:,,,,',,'1

the second fibre.

Main crack meets the secoc~ fibre and debonain5 p~oceeds.

Matrix crack froe t~e third

r.:: ":':':';';::-.-::_.1 ~~~~e c;:c~:~n!1ected to the

60

TABLE 9: Tensile Properties of Commercially Available GFRPE!

Property Test Method SI unit AV2 350 S AV2 310 .. ARNITE ISO DIN 33% GFR 33% Gf1I

TENSILE TEST

Tensile strength R 527 53455 MN/m2 165 185

Yield strength R 527 53455 MN/m 2

Elongation at yield R 527 53455 %

Tensile strength at break R521 53455 MN/m 2 165 185

Elongation at break R527 53455 % 2 2

Tensile modulus R527 53477 MN/m' 11 500 12 000

Typical properties of + I Beetle r PET

10% 30% 20%+

Property Test Method Units PET102F PET106F PETB04F

Tensile strength ISO R527 MPa 79 109 98

Tensile modulus ISO R527 GPa 8.2 11.4 9.5

Elongation ISO R527 • 2.3 2.5 2.4 ,

+ fl2.!D.e Retardant

30'%GFR . fLAME 45%GFR 55%GFR Properties ASTM UNITS RYNITE'V RETARDANT RYNITE RYNITE

530 FR 530 545 555 Tensile strength -40"C 0636 MPa 218 207 242 221

Testing speed 23"C 158 152 193 196

5mm/min 93"C 83 86 92 96

150"C 55 56 67 71

Elongation -40"C 0638 3 2 2

23"C 3 2 2 2

93"C 6 5 5 4

150"C 7 7 6 5

-I ~~~~ AI<zo ,,~

i ~~~of gll' ~ L,~

V ~~~1 E:i oW PIMt J.4- rJ~ """ cA. Co.) I/\.C..

61

Due to the relatively low strength and the ductility of the

matrix the major contribution to fracture stresses of the composite,

0c' are the stresses of the fibre, of. The fibre stresses at failure

are shown below:-

T-crack 90° of : 191 .3 MPa

45° of : 93.7

L-crack 0° of : 68.5

Assuming Hookes Law the tensile moduli are

(E )900 : 31 .02 GPa

(E)45° : 13.71

(E) 00 : 13.69

In the case of the T-crack specimen the mode of final mechanical

failure is catastrophic. Initially the crack grows discontinously, with

various micro-cracks growing at the tips of fibres, to distribute the

applied stresses at this point, where the stress is concentrated. These

microcracks compete to join the main crack which follows the weakest

path.

Eventually as the crack grows 'in front of the main crack tip in­

stantaneous failure of the specimen occurs. The stress concentration

at the main crack tip reaches a value sufficient to cause fibre pull-out,

fibre fracture and matrix failure simultaneously.

In the case of an L-crack, the mode of breakdown is dominated by the

behaviour of the matrix; crack acceleration can occur.

For constant strain-rate testing of specimens undergoing

L-crack it is possible to compute the magnitudes of the several

contributions to the work of fracture, W (in unit of energy (Nm)

per unit area of fracture).

1. For fibre fracture, Wf

usingwf = TTd'a' 1 N f P f

Wf = 5.83 Nm/m' (J per sq. metre of crack)

2. For debonding energy, Wd

Using Wd

Wd

3. For fibre

Using W P

W p

=

=

TTd'a' 1 Np f P

74.7 (J per m' )

pull-out energy, W

= TT d' a 1 N f P P 24

= 4713 (J per m')

p

4. For matrix fracture energy, -'VI .m

Wm = 6.22 (J per m')

Using these equations and the parameters in Table 10 Lhymn

and Schultz (19) calculated that the dominant process for L-crack

specimens is fibre pull-out. Our calculations although not agreeing

with Lhymn and Schultz show that this process is dominant. The

relationship between the critical stress intensity factor, K , and c

the total work of fracture, W, is:

where E is the composite modulus (31 Opal.

The existance of "boundary layers" (46) of aligned fibres

parallel to the mould wall leads to anisotropic properties. It can

be seen (Table 11) that as the plaque thickness increases the tensile

modulus (Et) decreases, since the higher modulus boundary·layers (in,

the direction of test) occupy a smaller proportion of the net cross-

section. When the sections are >6.4mm thick the boundary layer reaches

a stable thickness.

64

TABLE 10: PARAMETERS IN ·FRACTURE EQUATIONS

Parameter Definition

d

erf

lp

Ef

Nf

Id

Np

G"m

Em

Vf

Fibre diameter

Composite stress at fibre failure

Maximum pulled-out length of fibre

Fibre modulus

Number of fibres fractured per unit area

Debonding stress;tr:<I:, d f

Average debonded length same as average pull-out length

Number of fibres pulled-out per unit area

Matrix stress at failure

Matrix strain at failure

Volume fraction fibres

65

12~ Direct microscopic measurement

191 .3MPa Tensile test

Direct microscopic 500~m measurement

72.4GPa (19)

Direct microscopic 306mm-2 . measurement

191 .3MPa Tensile test

200~m

Direct microscopic measurement

-2 26l4mm Direct microscopic measurement

62.5MPa ( 19)

0.0061 Tensile test

0.33 From calculation

TABU 11 Tensile and Flexural Modulus results for samples tested: ~ the x direction (46) • see figure below:

"Topl1

Plaque thickness

(mm)

3.2

6.1

11.8

)'1 Hould fill/ direction

thickne{:;

::t

(GPa)

13.4

11.1

7.9

E;r

(GPa)

16.1

14.4

13.6

I 254mr:I

, /

surface 1 ~ iI'

i

(After ~etherhold)

GC>

The thickness effect on the flexural modulus was similar to that

of the tensile modulus. However since flexural modulus is a function

of the product of the area of the aligned fibres and the cube of the

distance from the neutral surface, the flexural modulus decreases

rapidly. There are considerations that have to be taken into account

when designing for strength with PET and glass fibres (illustrated here

using PBT). Tensile 7. bars and plaq"es of GFPBT were tested, on a con-

stant extension rate testing machine and results of tensile stress and

tensile strain at failure recorded (46). The mean values for five

specimens are shown in Table 12.

Table 12: Failure data obtained at constant extension

rate of 5 mm/min (Test temperature 23°C).

Specimen Tensile Stress Tensile Strain , (MN/m ) (%)

GFPBT ASTM bar 140.4 2.95

GFPBT plaque 0° 100.7 3.39

GFPBT plaque 90° 97.9 2.67

It is clear from these results (32) that, although the moulded

ASTM bar gives the highest failure stress, the highest strain to

failure is achieved by the plaque 0° specimen.

Due to the uncertainties of predicting loads to failure of

components moulded in fibre reinforced thermoplastics it has been

suggested that lower bound data be used where design for strength is

concerned.

67

Predicted and experimental load results are shown in Fig 15.

The results predicted from finite element analysis were obtained

using tensile moduli appropriate to the unfilled PBT and to various

fibre orientation distributions associated with the glass-filled

material.

The comparison between experiment and predicted results for the

unfilled PBT demonstrates the accuracy of the finite element stress

analysis. It is evident from the results that close agreement between

prediction and experiment is obtained with the unfilled component.

The results of glass-filled PBT show that if 'upper bound' modulus

data appropriate to the highly aligned FOD are used in the finite

element predictions, a considerable under-estimate of deflection for

a given load would result. If modulus data appropriate to the 'lower

bound' are used, the experimental results fall closer to the lower

bound. The predicted behaviour assuming moduli appropriate to the

random-in-plane fibre orientation distribution gives close agreement

with the experiment.

It has been shown that the matrix(Fig 16)and the composite(Fig 17)

show a much greater sensitivity to temperature than to strain rate (49).

The matrix material shows brittle behaviour from -SooC to room

temperature. At 60°C, a reduction in initial modulus and the onset of

plastic behaviour are observed. These tendencies are amplified at 120°C. bet ........ 1JJ ~ ""d.

The change in behaviouri· 60°C is expected as PET generally exhibits

a glass transition near this temperature.

68

o L------.-------r------~----_,----~ 0.2 0.4 0.6 0.3 1.C

!ieflection (mm)

r~gure 15 100 second isochronous load-deflection behaviour oor polybutylene terephthalate (PBT) and glass-fibre-reinforced poly­butylene terephthal~te (GFPBT) at 23°C, loading direction ''-'. (32)

(After Darlington" atICI iJl'ptrlo",)

LOAD (N)

500

300

.-80"[

.200 1 . RT

100l toe I ' i I If

.. 60·(

.120"[

PET - I

oL-______________ =-__ ~~--~----~--~ o ~ 8 12 16 20 2< 28

CRC3:;lEAD DISPLACEMENT (mm)

Figure 16 Load-displacement curves ·for the unfilled mat~ix (49)

(After Schlll tz 0."..01 F~)

6ao

'00

LOAD (N)

ino

530-1

o la J I

CR05.-;HEAD DISPLACEMENT (mm)

Figure 17 Load-displacement curves for the PET/glass composite (49).

(After Schultz <Wl ~dtich)

7()

For the composite)serrated force displacement curves are seen at

lower temperatures (Fig 1~. That is, regions of increasing crosshead

displacement with little change in load followed by sudden drops in

the load. At room temperature these easy deformation and load drop

transitions are smoothed out. As the temperature is further increased,

the modulus decreases and the load-displacement curves become smoother.

Fig 18 shows the mechanical properties for the matrix, where the

most striking features are the normalised strength, er T*' strength

maximum at 60 0 G and the tendency of room temperature properties to

change with decreasing deformation rate toward those observed at the

higher temperatures. For the composite, the expected drops in modulus

and tensile stress as the matrix goes through its glass transition are

evident (Fig 19} The greater matrix ductility is reflected also in

the large increase in the work of failure. The most surprising effect,

however, is the ultimate stress or work of failure both are slightly

higher at -80 oG than at -20o G.

To see the effect of fibres on the strength of the material all

the values were related to the matrix at the standard condition, Fig 20.

Three effects are wort~ of note:

1 . the principal: effect of the fibres is to maintain a high

strength at low temperatures, even while the matrix becomes

embrittled.

2. the peak in matrix strength moves towards higher temperatures

at higher deformation rate.

71

"'. ~

W·

rr • T

2

. 80 . o 8 20" · " ~ 3 RT ~~ RT "

-.---.-'~50

0 ,

" .12C • 0 , ! '------120

W,[ -- .. ~ 50 , 10,

" RT left ~ - ~ ·80

-..,----.:.

1~,1 -C)- ~ ·20

0

• 50 •

2 +

" ~'80 RT

" " " 1 ' 120 0 . 0 "80'(

o· 20'( " RI

0' : 11~:t

10" 10" 10" 100 10' 10' 10'

?ig. 18 Normalized mechanical properties for the unfilled matrix material (49)

(After Schultz ~ fr'tllcJ.Mh)

E'

0:' T

72

U

1.2

1.0

0.8

0.6

0.1.

~ 0-20 ~~

o RT e Gc:::::O

_~_--~;---"~o 120 0.2 s

2.0

l8

1.6

1.1.

1.2

lO

o o o

_----120

~ 0 0 60

~.:. , '·8~ ~: 0_ ~T' "

0.8 ~ ____

0.6

0.1.

0.2

-20 .

o

O~--------~ __ ___ 1.2

1.0 -80 ~ ~ ~~o-~i __ _ R' _o~'-

.. ~o ._' .• /~60 0.8

-20' ~.,-120 0.6 __ " . __ 0

0.2

---tl • • ea-[ o ·zo·c ~ Ri - E~r( o 1ZC·C

O~~------------~~ 10" la" la" 100 10' 10'

-1 CROSSHE~D SPEED (mm min )

Fig, 19 Normalized mechanical properties for the PET/glass composite (49).

(After Schultz·~

f'ru1l1.Ji.,0\. )

UNFillEC PET HATRIX

OL--_~9n~.--_~.n~~_~7--_~m~~G--~m~-~~~w~~~~~~~~'~~ TEMPERATUSE (oC)

Figure 20 Average normalized te~sile strengthO=T" (related to

strength of unfilled matrix at RT and 5 mm min- 1). (~9)

(After Schultz' and, f n:edMA-)

200,----------,

VISCOUS MAtRIX. fiBRE PUll-mII

100

xxxxxxxxxxxxxxxxxxxxxxxxxxx . (RA/lNi ANn (RA (KING . 'j,'j,'j,·j:J.1 Or Al fIB" ENOS 'j,'j,'j,'j,'j,'j.,'j,

'j,'j,'j,'j,'j,'j,'j,'1.'1.'1.'i.'I.'j,'j,'1.

BRlfllE HAIRIX .FlBRE PUll·OUT

-I 00 L,.-""",~-"-":""---:7;,---",,,-c:!. IO·J ID'! IQ'I IOU 10' I()l 10"

CROSSHEAD SPEED (mm min-1)

Figure 21 Failure behaviour aap for PET/glass composite. (~9)

(After Schul tz - cW1 &dIi.ch)

73

3. above 60°C, the curves for unfilled and filled material

approach each other.

For the composite, three different modes of failure have been

observed, (Fig 21). At low temperatures, fibre debonding and pull-out

are accompanied by brittle matrix failure. At intermediate temperatures

matrix cracks propagate through crazes formed at fibre tips. At higher

temperatures the matrix fails in a viscous manner, accompanied by

debonding.

1.4.1.2. TOUGHNESS

A composite must be, from a practical point of view, reasonably

tolerant to impact loading ie it must be capable of being damaged with­

out undergoing complete failure. Toughness is greatest when the,

length of the fibres is equal to the critical length lc (13). So

maximum strength and toughness cannot be achieved simultaneously, and

composites must be designed for optimum combination of the desired

mechanical properties. Fibres shorter than lc will be pulled from the

matrix rather than broken, when a crack passes through the composite,

this means that energy will be absorbed and toughness increased(F1j 22).

Two other methods are:-

i. The use of intrinsically toughened matrices (eg rubber modified

polymers) .

74

ii. The application of a soft coating to the fibres will act as

an inter-layer after the composite is fabricated. This has

been shown to reduce significantly the stress concentrating

effect of the fibres, especially under transverse loading.

Fracture toughness, Kc, can be def'ined (43) as the value of the

stress intensity factor Kr at which a crack in the specimen begins. to

grow unstable. This occurs at a load, F , taken as the maximum load c

in the cycle of crack growth.

Kc Fc Yft B=· specimen thickness = B.W,!

W= specimen width

a= crack length

Y= geometriC correction factor

The fibre alignment yields highly anisotropic properties, so it is

necessary to test in at least two directions.

a. with the initial notch and crack growth normal to the dominant

fibre orientation at the surface (T-cracks)

b. with the crack direction more nearly parallel to the fibres (L-cracks)

Specimens with cracks growing normal to the material flow front give

the highest Kc values, whilst crack growth more nearly parallel to the

fibres yields much lower values of Kc (50).

75

WORK OF

FRACTURE

w

1 .1

le FIBRE LENGTH "1.

Figure 22 Predicted depen~ence of composite work of fracture on fibre length. (13)

(After Folkes)

It has been found by various authors that there is a continuously

increasing relationship between Kc and glass content up to nearly

50% by weight of fibres. This tendency was observed for short fibre

reinforced PET, .by Friedrich, at least in the range between 15 and 55

weight % of fibres.

Increasing fracture toughness can be attributed to the additional

action of the energy-~bsorbing mechanisms that are triggered off by the

fibres in the material. In specimens with T-cracks, a high proportion

of fibres are orientated transverse to the direction of loading. This

leads to many specimens possessing a higher fracture toughness at a given

fibre loading. Kc data for constant strain-rate loading are obtained

from the maxima in load-displacement curves of notched compact tension

specimens; tests on Rynite 545 (45% by wt fibres) yield the following

results (19):

(Kc) Lexper 1

= 14.5 MPa m~

(Kc) exper = 9.9

45°

(Kc)T exper 7.7 =

Strain at composite failure is determined to be:

(E)L = 0.0061

(E) = 0.0068 45°

(E)T = 0.0050

For high fibre content the values seem to reach a maximum in

fracture toughness. After this lev~ it is assumed that Kc decreases

77

very rapidly due to lack of matrix material. In addition to the fibre

fraction two other microstructural parameters and their influence on

the fracture toughness are important: the effect of matrix ductility

and the modification of the fibre-matrix interface.

When two matrices of different toughness were tested it was found (47)

that the toughness properties of the composite in the range of low fibre

contents were not significantly improved c .···c when the

fibres were incorporated. With increasing volume fraction, on the other

hand, cracking is high~influenced by the bond qual; ·.ty of the fibre­

matrix interface, (Fig 23h which becomes a dominant element of the micro­

structure of the composite. It was shown that with a laboratory blend

with poor adhesion between glass fibres and matrix (lack of adhesion

promoter), the fracture toughness values dropped to about 70% of the

fracture toughness of a commercial product with the same fibre content.

If the properties of the fibre/matrix interfacial layer are poor, the

cracks tend to propagate exclusively along these regions. ThUS, their

fracture properties determine mainly the properties of the complete comp-

osite.

Owen and Bishop (51)carried out fracture tests on a thermosetting

polyester resin (having therefore very different matrix properties),

containing various forms of glass reinforcement to investigate the effect

of stress concentrators on the failure of reinforced plastics. Onset

of crack propagation occurred when the elastic stress distribution

around the crack tip reached a critical level characterised byt he

critical stress intensity factor Kc. For brittle materials Kc was found

to be independent of crack length and is regarded therefore as a material

constant.

1

Kc( E?a "'-~"l

� 1, __ .-__ .-_---;,-_---:

" 10

. sr---t----2~r---j--;.,----j

-I

-,

~ ... : -. , ; .. .

-B):n.tp' ·0 0 ~otnx I •• IiJfrl~ 11

D·~0--~IO~~I~a~,~u--4~C---~~--~'Or-~1.~OO

w/~,(,. fibres

Figure 23 Fracture tough_ ness K ot-short fibre reinfoFced P.E.T. depending on fibre fraction and matrix toughness (matrix II matrix 1) (47)

(After Friedrich)

The applicability of linear elastic fracture mechanics to glass

reinforced polyesters has been found to depend on the type of re-

inforcement and the position of the crack. With the crack parallel

to the main body of fibres Kc is found to be constant over the range

of crack lengths, Kc is small and relatively constant for L-cracks.

The fracture toughness of a composite material is a very

important engineering property. The area under the stress-strain

curve up to the point of failure, is a measure of the work fracture.

This is formally related to the fracture mechanics parameters G~ and

K: _, (critical strain energy release rate and fracture toughness resp­. le

ectively). As with many composites, the conditions that lead to high

stiffness and strength also result in low elongation to -failure, so that

work of fracture can be very low compared to that of the parent matrix.

If inherently brittle fibres are added to an otherwise ductile

matrix, the impact strength of the composite decreases rapidly as the

fibre concentration increases. This effect occurs unless the ductility

of the matrix is suppressed, as will take place at low temperature. In

this case it appears that the addition of fibres lead to an increase in

impact strength.

Very little has been written about the impact behaviour of glass

fibre-reinforced PET, although some tests have been carried out by

companies on their own products, Table 13.

The most recent work by Friedrich (52) shows the current level

of understanding of fracture toughness of the thermoplastic matrix, this

is affected by the volume fraction, orientation and distribution

TABL:: 13: lJ:'ougnness of Commercially avail!} 01e G~PET

Property

Tmpact Test

impact strength (Charpy); unnotched notched impact strength CEAR?Y IZC~

Hardness

5all indentation hardness, "358/30 H961130

Rockwell hardness, scale L sC:;tle ., scale .:.

Sho~e ~2rd~e3s, scale D

Abrasicri ~r:'!-method (emery -cloth) Taber CS17, 10 ~ load

Property ASTE

::i:zoci impact strc!lG'th _40°C D256 (notched) "" ... 0,... c:; t ..

:::ocl-::well ha::-d-neS3 D785

r-:-opert:;

~so

2 179

R 179 ~ 180 "

R868

ur~ITS

J!m

Charpy i~pact strength liotched linnotched

Izo~ impact strength notched Tj:1:1otc~e~

DIl{ f...STH

53453

53453 D

53456 53456

D D D

5.3505

53516 D

530

9(; 101

~j100

:~ 120

. T' 2 l:l.</ s? kJ/m~

J/!':l v /;.,

81

~ kJ/m-

kJ!:n 2

25'0 J/r.:

.... 2 j'" '/"'2 ~ilVm

785 785 785

mm 3 10!.:-l;. "';;/

100e.

RYNIT:: F? 530

80 85

H100 ~i20

7.0 22.0

33 9~

rev

545

123 117

1:.00 ~~20

7.8 24.7

30 162

25

.., -( .;!

~O

240

114 102 61

8c

240

45

~-~ /).'

123

AV2 370

304:

18.5

5:: 30~

35

10 0_

.'))

230

, 15 103 64

90

200

38

of short glass fibres and their interfacial bond quality.

An increase in the thermoplastics toughness can be expected

with increasing extent of reinforcement if the matrix is in a brittle

condition and if the fibres are well bonded and mostly oriented

perpendicular to the crack front. An opposite tendency may occur

for matrices which are ductile, in the presence of the fibres.

The trend of the variation in a composites toughness can be

described as an empirical relationship.

Kcc = M.Kcm

where Kcm = matrix toughness

and M is a "microstructural efficiency factor"

The magnitude of M depends on volume fraction and FOD over the

cross-section fractured. M can be written, therefore, as M = a. + .1t.R

which leads to the "reinforcing effectiveness parameter" R being

directly related to Vf and the geometric arrangement of fibres.

This is naturally different in L-as compared to the T-direction.

Term "a" is a "matrix stress condition factor" which reflects

changes in fracture toughness of the matrix as a function of plaque

thickness and/or the presence of fibres. Normally "a" should be equal

to 1. Generally this approach can be used to illustrate how certain

reinforcements (R) can influence the fracture toughness of various

thermoplastic matrices, and how in a given fibre/matrix system the

toughness can be systematically optimized, Fig 24.

82

3 AKcco' i / ~ ~1const / /

~ n=cor~st / 2~i /f

/1 o ____ 12 2 /

/ 1

/

/

/ n > 0

decreasing temperature

fibre strength aspect ratio increase with

o #

~--------~'----~----~--~~--~R ~ poor ~nter- J I

'- face . . ..... ......... l.ncreae~ng

matrix ductility

-1 1 ",.

"l-L ____ : I ~ -2 '-....

-3 J

I

1 Roconst

.6K~c1 / r; 1

~ 2 n=const

,-" 4

n<o

Figure 24 Possible steps of toughness improve­ment (AK' t) by manipulating factors Rand n (52). cc

(After Friedrich)

83

-% of fibres transve rs e to crack

fibr!: o1'"ie!!.tc.­ticr:

- fibre conte:::.t

1 .4.1'.3. STIFFNESS

The stiffness of a composite is given by:-

E = no v Ef + (l-v) Em .. - -. - ~ (I~}

where n = orientation efficiency factor 0

Ef = Young's modulus of the fibre

v = Volume fraction of fibres

E = Young's modulus of the matrix m

If the fibres are not all the same length then either the distribution

of fibre lengths can. be represented by a single number average fibre length

or nl must be obtained by summing the effects over the entire population

of fibres.

The stiffness of short reinforced thermoplastics depends on the

fibre length (and/or distribution), volume fraction of fibres, the stress

transfer efficiency of the interface and the fibre orientation (53).

Figure 25 shows very clearly that the stiffness is a very sensitive

function of angle for small degrees of off-axis loading. The consequence

of this is that unless the fibres are perfectly aligned, the observed

longitudinal stiffness will be substantially less than predicted and the

difference in composite stiffness between using eg carbon or glass as

the reinforcement will be small. In terms of the commercial exploitation

84

TENSILE

HODULUS

IN 106 PSI

1.3~--------------,

1.1

1.0 •

_____ TyPE I BARS

SHAFFER MODEL LEES MOO~L

o 20 40 GO 80

ORIENTA TIOll IN DEGREES FROH STRESS DIRECTION

Figure 25 Angular dependence of modulus for 30% glass-reinforced PS']' (53).

(A fter Hciially)

of high modulus fibres in thermoplastics, these observations suggest

that there is no advantage in using any high performance fibre other

than the cheapest eg glass. However real components are subjected to

multiaxial loading, where a perfectly aligned short fibre composite

would be inappropriate.

The introduction of glass fibres increases the average values of

stiffness and strength over those of the thermoplastic matrix (38), but

since the reinforcing fibres are distributed and partially aligned

by the flow of polymer-fibre 'melt' during processing, injection moulded

articles fabricated from fibre reinforced materials often exhibit an

overall mechanical anisotropy.

Coupled glass-fibre reinforced polyethylene terephthalate is an

injection mouldable material which is available in a number of different

grades. The grades can differ in the method chosen for chemical coupling,

in the amount of glass fibre present and in the compounding method used.

The distinction between GFRPET and the uncoupled· material is important.

Polyethylene terephthalate reinforced with glass fibres is known

to be among the construction materials with the highest stiffness values.

Increasing the thickness (54) of a GRP moulding is more likely to lead

to a high modulus in the direction perpendicular to the major flow

direction since the thickness of the surface layers remains constant and

this region determines the degree of anisotropy and the direction of

greater stiffness. The balance of fibre orientation in the core and

surface layers is of considerable importance when comparing flexural

stiffness and tensile stiffness and when dealing with the strength of

these materials. Table 14 shows the stiffness properties of some

86

TABLE 14: Stiffness properties of 'commercially available GRPET

Property

Flexural test flexural strength flexural stress at maximum load flexural modulus

Creep test (1000 h) creep modulus at 1~~ strain stress at 1% strain

Property

ISO

178

178 178

Flexural strength Fle~ural modulus

Property

?lexu:-al modulus

Test He thad

53452

534-44 5344-4

SI UHIT-

~

w'/ C. 1'~11 In

I 3EETLE 1 PET

_40°C ?_o,... ~, v

930 e 1500 e

ISO 178 I30 17G

A.STH

;)790

D790

U1~ITS

HFa GPa

U!:I'~'S

GPa

HPa

07 " ,

530

10,3 9.0 3.6 230

ARl-: ITE A (?3T) .~V2 360S

250

-:~ 500

102Y

148 7.7

106?

155 "':0.1

lENITE F'''530

11.0 10.3 4.3 220

AV2 370

280

12 000

8000 80

545

15.2 15.8

5.::> ;>;<-_.;)

153 ?7

555

20.7 0'" 0 , , . , ':.2 310

commercially available glass fibre reinforced PET grades.

For predicting the tensile stiffness behaviour of a component

highly accurate creep machines are needed and the procedures recom­

mended in BS 4618 Part 1 should be followed.

The results for GRPBT are shown in Table 15 (32).

Table 15: lOOs creep modulus data (lOOs, strain 0.1%)

Material Specimen Tension modulus

(GN/m')

PBT Plaque 90 0 2.5

GFPBT ISO II bar 10.0

GFPBT Plaque 90 0 5.0

GFPBT Random-in-plane 7.0·· .

For glass-fibre reinforced PBT, studies have shown that the data

obtained on plaque 90 0 specimens are a reasonable representation of the

lower-bound behaviour.

1.4.1.4. ENVIRONMENT

In the last few years there has been an increasing awareness of the

effect of the environment on the properties of polymer based composite

materials. Generally the corrosion resistance of fibre reinforced

plastics is superior to that of metals or alloys.

88

It is well known however (55) that reinforced plastics can be

weakened under an aggressive environment. In composites the weak

feature is usually either a glass fibre or an interface. In this

context the polymeric matrix protects the composite. The environmental

durability of a composite is closely related to the susceptibility of

glass fibres, or the matrix, to a specific environment. In most

environments GRP is reasonably inert, especially when the component

is not subjected to service loads, Recently, however, it has been

shown that even in an aqueous enviornment in the presence of a

sustained load stress corrosion cracking can occur.

A disadvantage of thermoplastic polyesters is their indifferent

hydrolysis resistance: Prolonged contact with water at 95°C, or even

as low a temperature as 50°C has significant detrimental effect on

properties (56). The deterioration in mechanical properties caused

by hydrolysis occurs rapidly at the higher temperatures and relative

humidities and progressively slows as the temperature and/or humidity

are decreased (57).

The uptake of water in some polyester glass fibre composites at

temperatures as low as 30 0 C can lead to blister f.ormation and permanent

microstructural damage which has a large effect on the mechanical prop­

erties (58).

PET is attractive to moisture and can hydrolyse rapidly when melt

processed or if exposed to high-humidity environment subsequently.

Melt degradation is further complicated through simultaneous thermal

and thermo-oxidative effects (4,5). Zimmerman (4) has presented a con­

cise summary of the complex inter-relationships of degradation mechanism

of PET.

89

On long term humid ageing in hot water, impact behaviour especially

is rendered more complex by simultaneous crystallisation, molecular re-

ordering _ and losses of interfacial bond strength. The hydrolysis of

PET, over a wide range of conditions in the solid phase is known to be

auto-catalytic. At 85-87°C (20) reaction rates are independent of thermal

history (and therefore independent of microstructure and/or crystallinity)

since-, moisture transport rates exceed those required to hydrolyse the

polyester chain.

The constant deformation rate tensile data in Fig 26 convey the

dependence of composite fracture strength on hydrolysis time for specimens

of various (initial) molecular weights. Number average molecular weight

decreases more rapidly during solid-state hydrolytic depolymerisation for

compounds containing high volume fractions of glass fibre.

Losses of mechanical strength and toughness are apparent only below

critical molecular weight levels: retention of tensile fracture strength

is determined by fibre volume fraction, whilst pendulum impact strength

is modified substantially only when matrix embrittlement is reflected

by losses of crack initiation resistance.

Losses in interfacial shear strength through physico-chemical

attack of glass-resin coupling reduced the work of fracture for debonding

and subsequent low-friction pull-out.

o.nd Sc.lwl rz. Lhynni(59) studied the effect of salt solution on PET/glass composites,

and found that a neutral 10 per cent Nacl solution had no effect on

fracture strength. The composite structure is unaffected by treatment

for 12 hours. Methanol, also, was found to have no effect on fracture

strength.

90

'\ (~1I nit c Str(~lIgth

UUlm-2

)

F igul"C! cl6 -

IW

2U

2

A(.'(.',elt'rlltcU 1 . tellsi."le st.· l(,.\~ rt~(·tI\O . ,lcllgth. (20) (> ill \.)!ltl'r 87

0

(1'., "" I ' (AfiAr Ha~ocll)C: e{(ect of ' ini\:'iol molecular wc' ight 011

l.:lIId s '\r " of ilgC! ing / .!. 27.) Une • )

I1gh levcls . typicatJ Dllt)' nt cx~r~m:leyxlt~~"lI{'lY 5m.,11 (

n:T-A

Illitinl fill

I~~·) o 34,5W 6 lU,6l10 0 16,8UO 0 16,{)(1U

o

4 6

alld may leat! to StlltlStic.:Jl

mensured

illsip,llif' 1 lC [(nce

I'ET-U

Initial HII (r, 11101- 1)

• 34,50U • 17,600 T 15,800 • 15,00U

.-

The lifetime behaviour of samples exposed to these environments

was similar to air exposure.

A very important factor relating to environmental effects in

composites is the combined influence of stress and environment. In

stress corrosion, the lifetime under simultaneous chemical and mechanical

influences is reduced relative to either influence acting separately.

Environmentally-induced degradation of materials at the tip of a

crack enables the crack to propagate at a relative low level of applied

stress. At the same time the propagating crack continuously exposes

undegr"ded' -- material to the action of the environment.

The degradativ.e,' attack can apply to the matri~ ~to the fibres, or

to the matrix/fibre interface. If it is fibres that are principally

attacked, then the matrix can act as a protection. However this protec­

tion is interrupted by local mechanical failure.

It has been shown that glass fibres crack spontaneously in acids,

even in the absence of mechanical loads. The design of many GRP'sis

limited by the loss in integrity which occurs with the onset of micro­

cracking at low composite strains (60), occurring particularly in those

areas where the principal fibre direction is perpendicular to the

direction of an applied stress. Loss of composite integrity may be

important in applications where any flaws or microcracks may expose the

fibres to corrosive attack.

Weakening of the glass fibres has been attributed to the ion

exchange between surface sodium ions of the glass and hydrogen ions from

92

the acid, and it was claimed that the volume change on hydrolysis

produced surface tensile stresses (61). The subsequent shrinkage

of the outer layer of the fibre results in surface tensile stresses

which lead to failure. A second explanation for fibre cracking is

based on the leaching of the material at the tip of flaws and in

changes in the surface tension of crack surfaces.

A third .explanation (55) is that weakening of the fibres in an

acid environment takes place by removal of aluminium and calcium

elements from the fibre. The local depletion of aluminium and calcium

could cause a local pitting which acts as a stress raiser for the

eventual fibre cracking(see Fig 27 and Fig 28J

The removal of calcium and aluminium in a fibre produces an open

structure which is less retardant to chemical attack than the initial

structure. Perhaps more importantly, the local pitting produces stress

raisers. Bond scission is ea5ay accomplished at a region of local stress

enhancement during slow propagation of the crack front. The end

result is the microcracking of a fibre at various locations along its

length, Fig 29.

Fibres are easily fractured around the crack tip due to the combined

action of stress concentration and chemical attack (62). Easy fibre.,'

fracture on the crack front means that fibre pull-out does not occur

and therefore the fracture plane is quite flat (63). The implication

of the elimination of the fibre pull-out process for fracture is that

the energy of fracture is reduced, eventually lowering the fracture

toughness (64).

93

P. 5 L I

02

C R

2

«J'SOP ':t'.EV ,:e2. 72e E[IAX E [lA X

:~g 27 EDAX profile of cor­roded specimen (55)

(After Lhymn & Schultz)

( . ) , , fibre

T

c!"aCK o~eninb

filled ":i th /1' r,o"..utior/

(~) ~ : I

CUPSOP 'kEv,:e2. 729

Yig 28 ZDF.X profile of uncor­roded (&ir) specimen (55)

(After Lhymn & Schultz)

2.-:: -:.:;":: ::=os..~::: ""..:l'-;-; ur.de~· a.r: ag~:'e5c~vs env~;on~e~~ (55)

Grack reeches 8 fibre

04 ..

In an injection-moulded polyethylene terephthalate matrix

reinforced with short E-glass fibres, the environmental durability

depends critically on the kinds of mechano chemical attack occurring.

In alkaline solution the PET matrix and the glass matrix interface

are weakened whilst the fibres remain unaffected .

. From Table 16 it can be seen that the glass fibre deteriorates by

a chemical depletion of Ca and Al in acid environments whereas the

fibre remains quite inert in neutral or basic environments. Attack on

the PET matrix occurs with 10% NaOH soln where extensive matrix pitting

occurs and also partial debonding (65). Alkali particularly attacks the

rubber-like particles that are added to commercial PET/glass composites

leaving profuse pits (former sites of rubber-like particles).

Friedrich (62) performed stress rupture tests with Rynite 545 with

fracture in the L-direction. The results for these tests performed in

air, water and three acid environments are given in Fig. 30. These

results show three modes of failure.

The composite is relatively unaffected by testing in water and the

data points for fracture toughness follow very closely those determined

in air. These specimens fail by normal crack propagation under the

plastic response of the material to the applied load without being

affected by the environment.

In acid environments it can be seen that there is a reduction in the

initial fracture toughness necessary to induce crack growth and ultim­

ately final fracture. It can be seen that in the acid environments the

95

?eaks (55).

:::nvironment 1 1 1 , '- , '-!:,ledium Al/ Si '""a/ "i

.""'I.J..r ~ .,-v_'-) O.4G

1 J;.:; I-laOE 0.23 0.50

1 rv': v.- RCI 0.07 0.12

1 cr,:: H -0 -·2/:) 4 0.15 0.30

1"-': v .. PT··O ,." 3 0.07 0.14

10;',. ;";a::: 0.23 0.50

j·ie t:lar.03.. J.22 0.50

(After L~ymn & Schultz)

10.-------,----.,.----------, Rynile 51,5

9 l_l-direClion

8 K --- 1I -1-11I-

.::,--5

---

-------- --2'~-----L------L-----~----~

-v pre-Ireoled 200h in 10% Hel ..... nc fraclure unlil Ihis lime

10' 105 10° d

• I ime of failure I o

'.:.'ine to ?2i2.u!"e ( nee

(~fte~ ?~iedrich)

97

)

~i~urc 30 i~itial 3t~es~ i~tensity-~i~e to failure curves fo!'" :tvr..ite -: 5h 5-1, L-~irectio~ ~; .. ~ V"".,.....; ("'·u~

~ --, --- .. -- ......... envirOllse~ts (62).

stress necessary is an indication of the aggressiveness of the

environment 10 vol% HCl and 3 vol % H2S0~ show a decrease in the

necessary stress level to initiate cracking.

Fig. 31 shows comprehensive data for hydrochloric acid tests

of Rynite 545 and Rynite 530. It can be seen that the initial

stress intensity factors, are higher for T-crack specimens and

higher for the higher fibre loadings. T-crack specimens however

experience a much steeper decrease in stress intensity factors nec­

essary to induce stress corrosion failure than L-crack specimens.

This decrease seems to be due to the higher amount of fibre fracture

usus ally responsible for the higher fracture toughness of T-crack

specimens has changed into fibre pull-out after degradation at the

interfaces.

It is important to note that when Friedrich (62) pretreated

specimens for 200 hours in 10 vol % HCl at room temperature without

exposure to an external stress and then loaded them, no significant

influence of the pretreatment could be observed in air or 10% HC1.

This shows that chemical attack during immersion does not normally

cause serious loss of physical properties, but that a combination of

chemical and mechanical interaction between the composite and the en­

vironment leads to rapid deterioration.

During the corrosion process the tensile stresses produced are

capable of developing to levels sufficiently high to bring about spon­

taneous cracking. No degradation effects on the thermoplastic·.matrix

by the acid solutions are observed.

.' ~ ~

V>

.so -

10r.:Kc----;-----.,.-----.,.-----,----...,

,/ :~I~-gr- -_

5

4

JI

I

530-1 [-direction

_Rynit o

in 10% HCt

5~5-1 l-dir~ction

530{l l-direct ion

2 '" prelreoted 20Gh in 10 '/, Het T I jm~ of ioiiure !

id

Time to railure

7i~u~e 3! I~itial stress intens~ty-tirne tc f~ilure c~~ve~ ~o:· tie ~.,- end ~­directionG i~ Ry~ite ~

i~ 5~~:v~~~~::~~O(~:. ~ ~t2. i~e d

( s~c)

As has already been stated alkali solutions attack the matrix

rather than the fibres (65). For L-oriented specimens, severe matrix

cracking and interfacial debonding are seen, fibre fracture being a

minor event. The matrix fracture initiates from either fibre ends or

matrix/fibre interfaces.

For the T-oriented specimens again the general failure proceeds

by interface and matrix fracture. First, discontinuous matrix cracks

are initiated either from ends or within the matrix itself. Later,

discrete matrix cracks join together. The average path of the matrix

cracks is not highly aligned with the mean general crack direction;

the individual matrix cracks follow paths dictated by the microstructural

detail.

In NaOH, rapid interface failure p~ovides an easy path for crack

propagation. The fact that the matrix fails without excessive deform­

ation means that:

a. the rapid interface failure has made failure of the matrix

possible without viscoelastic deformation and/or

b. the matrix has embrittled.

It appears that both phenomena are working together. In the

T-geometry, the crack propagates mainly through the matrix phase,

since the fibre axis is macroscopically parallel to the direction of

crack propagation. Thus, the presence of fibres does not cause any

obstacle to the crack growth behaviour.

The effect of the environment has been studied in great detail for

acid and alkaline solutions by Lhymn, Schultz and Friedrich (62,65).

100

They have even investigated the electric break-down of polyethylene

terephthalate glass composites (66). The work by these invest­

igators appears to represent the total knowledge of environmental

effects, apart from that shown in Table 17 which is a commercial

evaluation of PET/glass composites environmental resistance.

ARNlTE Table 17

Acetic acic. Acetic acid Acetic acid Acetone

1::,. /"

1 o;:~ 100){

AmmoniUE hydroxide 1~~ Ammonium hydroxide co~c. i;.r:ile

Benzene Bleaching lye 5i-ake fluid 5ut~ne

Butanol Butyl acetate

Calcium chloride 10% Calcium hypochlorite Carbon disulphide Carbon tetrachloride Chloroform Chromic acid Citric acid Cottonseed oil Cresol

ueterger.ts

Dibutyl phthalate Diesel oil Dioxane

Ethanol :sth~!'" (diet::yl-) ::'th:.~l acetate Zthylene dic~loride

:reo::. i .... ··

4o;:~ 1~,"

15;' 25%

--' 7;~

gf)'~:

~lyce~o~ (~lyce~1~~) Glycol Grea::e

;{CA~rrc

~ydr~chloric acid ~yd~oc~lo~ic ~c~c

E';drofluori: acid Hy6rofluo~ic acid Hydrogen peroxide

1 Cf.,; cone.

Che~ical resistance of ?ET/slass co~?osites

+ o

+

+ + o

+ + o +

+ + +

+ +

+ o

+

o

o +

o

o o

o o

o o

o +

o o o

o o

o o C'

unfilled c!"Ystalline

~_c~ v'Coc "oor '::::) '.... 0 ! (: '--'

+

+

o +

+

+

+ + +

o +

+ + + +

+

+ +

+

+ + + +

+

+ +

+

o

+ +

c

c o o

+

o o

+

+

o

+ o

o

o

" o

+

+ +

+

c

g:!.asG-:illed cr":tstnlline

?~o~ ·~Cc~ R~o~ -J v "' ,,,", L 10"

+

+ o o o

+

o

+ + o +

+

+ +

+ + +

+

" + +

+

o

+

" c

o c

+

c

o o o

o

+

o o

+ +

+

"

o o +

o

c o

o

c

o o o

+

+ +

o

i..rtNITE Checical resistance

:~e!"ose::e

Eeth~!lol

Het:::tylene chloride ~ethylethylketone

Hineral cils Motor oils

Eitric acid 1-!itric acid l":itric acid

Oleic acid Olive oil

Pe!"chlorethylene Petrol ?etroleu:c et"e:' Phenol Phosphoric acid ?hospno:-ic ?hosphoric acid Potassiuc c~loride

10;~ ',0;; 7-::;;( cone. )

100%

3;j 3~~ 855~( cone. ) 10;.;

Potassium dichromate10;~ Fotassiu:J h: . ."c.ro::.:.ide 1% Potassium hydroxide 1 a:,~ Po b3S iu!:! hydroxide607; Potassium per!!langa!1ate 1O:~

~ ilicone fluids Soap 6o~ution 30dium bicarbonate Sodiu~ bisulphite :iodium ca!"bonate 50diu:!! ca:::-"oona te ,odium chloride

1~~ 1O:~ 107: 10% 2O:~ 10;;

+

o

+

+

+

o o

+

o

+

+ o +

+

+

o 30dium hydroxide 1<~ +

lodium hydroxide 10,:: >odiuo hydroxide S~~ ;odiu!:! hypochlori te107j + ;ulphuric aC id 35'; -t-

;ulphuric acid 3(J';b + ;ulphuric acid 98%(conc.) -

:e trahydro fur"n ~oluene

~ran5former oil ?richlorethylene ~urpentine

'aseline 'estable oils

at er 'hi te 3pi:-i t

ylene

o +

+

+ +

+ .;-

o o

o o

o

o o

o

o

o o o o + o o o

+

o o

o

o o

o o

.;-

+

+

+

+

+ + + o + .;-

.;-

+ .;-

.;-

o

+

+

+ +

+ + .;-

+ + o

.;-

+ +

o +

o +

+

+ ~

tt!!filled

+ +

o

+ +

+ +

o

+

+

+

o

+ + +

+

+ +

+ +

+

+ +

+

+ o o

+

+ +

+

+ +

.§:lass-:illed c:'~;2.t?"l::'ine

,:::0,.... ,. :~O::·.. Q()o .... -_..... '., ~ "''- ' ..

+ +

+

+ + + o +

+

+

+

+ +

+ .;-

+ +

+ ... ...

o

+ o +

... +

+

o +

+

o

+

o

+ ...

o

+

o o

o

o + +

+

+

+

o

+ +

+ +

+

+

+

+ +

+

ARNI~ ~esistance to groups of che~icals at room tecperature

unfilled glass-fillec . Chenicals' arr:o:-phous cr-y.stalline c:-ystalli!le

Inorganic acids concentrated (non-oxidizin5) e.il u tee. ( 1 : ~ ) + +0

hig~11:t diluted + ~ + •

I:!o:"ga:1.ic acids concentrated (oxidizinc;) diluted ( 1 : 1 ) 0

highly diluted + + +

Org:u!ic E!.C icia concentre~ted 0 0 0

diluted ( 1 : ~ ) + +. + highly diluted + + +

3ases concentrated diluted (1 : 1 ) 0

highly diluted + +

3e..lt solutio:ls acetous + + .;-

neti.trhl + +

00.::lC 0 + 0

~liphz.tic tydr·:.-ca:-bons 0 + +

oils cJ::d o!'e3.s~.s + + +

P. romn. ~:"c ~1ydro-

carbons + 0

phenols

t;l,-.,'~ ha ... Oc,e __ a'tee. pe!'halo~;ena tee: + + + ::ydroca="8ons .... ~_.j...1 -.

='~- "'-,: haloge~tec.

.~.lcon:)ls monov.:::le!'lt + + .+ pol:r:ale!2 t ~ 0

r~e tone::; , 0

o

.:...":.hers

+ = resistant; no attack, no change or only a very slight change

in weight (less than 1%),

reduction in tensile strength at break remains under 10%.

o = partially resistant; in course of time there is a distinct

deterioration in tensile strength at break (10-50%) and a

change in weight of 1-5%; in many cases a short contact may

be considered permissible.

= non-resistant; after a short time the material is seriously

affected and/or dissolved; change in weight of more than 5%

and/or a reduction in tensile strength at break of more than

50%.

105

1.4.2. PETIPC BLENDS

There has been a considerable commercial interest in blends of

polyesters with bisphenol-A-polycarbonate. A number of workers have

been investigating these blends and they have been the object of

numerous patents (78, 79, 80, 81, 82, 83, 84, 85).

Chen and Birley investigated the properties of PBTIPC blends

(86, 87, 88) and followed up this work by investigations into PETIPC

blends (36). They demonstrated that performance of the blends was

superior when compared with pure PET and pure PC. Blends do not

suffer from shrinkage Qr become brittle during annealing, and remained

ductile although the 80/20 PETIPC blend exhibits increased yield stress

from 56.9 to 78.9 MPa. The difference is associated with the PC con­

stituent in the blends which appears to retain its softening character­

istics and thereby stabilise the dimensions and modify crystallite growth.

The 80/20 PETIPC blend also shows a comparatively low MFI.

A considerable number of polyesters are miscible with polycarbonate

(89). It was observed by Cruz et al (90) the PET was partially miscible

with polyca~bonate. Polymer incompatability is the general rule in

blending (91). This situation arises from the very small entropy gained ~

by mixing different species of macromolecules. Nevertheless, the

degree of incompatability varies widely and is of tremendous importance

to the morphology and to the ultimate mechanical properties of the blends.

PET cyrstallizes rather readily and it would be expected that PET

would crystallise from blends with PC. The extent of PET crystallinity

for blend samples after extrusion were assessed by thermal analyses (35).

Since PET crystallisation was quite limited in every case, further

106

crystallisation occurred upon heating in the DSC.

The results Murff et al (35) obtained for extruded and injection

moulded blends of PET/PC are shown in Fig. 32 and Fig. 33 respectively.

The sum of ~Hf and ~Hc is indicative of the level of crystallinity.

The level of crystallinity for extruded PET is about 15%. Injection

moulded PET has a similar level .of crystallinity up to about 80 w/w %

PET, contrary to the conclusions reached by Murff et aI, after this

level the crystallinity increases up to about 30% at 100% PET. The

reason given for this was the greater stress experienced during injection

moulding as compared to extrusion.

1.4.2.1. STRENGTH

The moduli and yields strengths of PET/PC blends (35) are shown in

Fig. 34. The ultimate stress at failure (Fig. 35) shows the decrease

in yield stress as the blend becomes richer in PET which has a lower

yield stress than PC.

The interesting feature was the elongation of failure (Fig. 36).

The blends failed at higher strains than the pure polymer, and blends

containing 60 to 80 w/w% PET did not fail within the available cross­

head movement (200%). The nature of the stress-strain diagram for

blends in this region are contrasted with pure PET in Fig. 37, showing

clearly the improvement over the pure PET.

After elongation the tested specimens were examined and it was

found that there was a large increase in the level of crystallinity

in the necked region as a result of the drawing process which occurs

(cig 32 Heats of f'usign, bHI ,and cry5te.l~iz.ation ,00He ~er extrusion (35)' ,

.. ~

(Aft~r Murf~ et all

'l 'T 5

-10 :---::!;o--:';;---:'::---,!;:---,,,} o 20 40 60 80 100

Fig 33 flea ts of fusion". tJ.i{ and crystall.i.·zsHon, tHe, aiter injed::ioo ftlolc'iing (35)

CAfter Murff et al)

'350,000 l-

Modulus (psi) f'\. r-

• • • • ~..-c-'

-

JOO,OOO L-_-'-_-L_--'-_-.-J'--.--I

Yield Strensth . (psi) .

10.000 ~

9aoo:..~ ___ -=:._

ecoo~ 7000!:-, ----:f.o---:'::---!:--~--! o 20 40 60

•

PC WEIGHT ~

Fig ,34 Modulus and yie-ld st.rength for inject.ton ~olded blends (35).

(After Murfl et al)

10'/,

(psi)

Stress

?ig 35 blends

12.000 .--.,--,---,-----,--,

10.000

8000

4000

2000

GL-~~~--~--~~ o 20 40 60 80 100

?C

~:ti~~~e strength for injection ~olded (35)

(After Eurf: et

• 100,

~O

~O

"0

lO~~~~~~~~~, o 20 40 60 80 100

blend3. :long~tion at break

Eaxi:.mm availabl<:! elongatior.. C.35)

for injectio~ ~Glcied cross~ead e~uiv~le~~

( p:; i ;

(~fter Murff et al)

------ __ 50 % P~T __ --------750,;. PET

°Or---~--~5~0~-~-~IO~O~-~-~I~sO~-~-~?~r~.r-'~ ::LO!:G;':. ':'ton

?i~ 37 Typic&l st~es5-strai~ dia~ra~3 i:l~=t~~tins ::ir;!": (;'ilctilit~.- of 60 anc 75:-~ ?~::' ble!1Qs (3~)

, " , ~"t. Cl.!..1

during mechanical testing. This is especially apparent in those bars

that did not fail.

1.4.2.2. TOUGHNESS

The toughness of a PET-based blend with PC is increased (92) due

to the outstanding impact resistance of PC. Bisphenol-A-polycarbonate

(PC) is an amorphous, high glass transition temperature (T 145°C) g

thermoplastic, characterised by an exceptional toughness. Engineering

thermoplastics (PC and PET) are classified as thermoplastic polymers

having reasonable-to-exceptional levels of toughness (93). There is

an increase in the impact strength of the blend with number of ex-

trusions.

1.4.2.3. STIFFNESS

PET is chosen as a blending component to obtain an increase in

stiffness. PET based blends offer an attractive stiffness over a wide

range of temperatures. PC and PET are thermoplastic polymers having

enhanced rigidity at elevated temperatures. This property is frequently

a consequence of a high. glass-transition temperature (T ), and many g

engineering thermoplastics contain aromatic ring structures of these

materials make melt processing and fabrication of parts more difficult

and lead to the necessity for definition of the melt viscosity behaviour

of these polymers at processing conditions.

110

1.4.2.4. ENVIRONMENTAL

The solvent resistance of amorphous PC is often very poor. Use

in some fields is limited by its low temperature behaviour, high melt

viscosity and comparatively poor hydrolysis resistance. PET is chosen

as a blend component to obtain a marked increase in the resistance to

chemicals and above all to fuels. The PET/PC blend has an inherent

chemical resistance, it is hoped that blending overcomes the limitations

of PC.

Some of the consequences of PET being incompati-ble with PC, at

least in the solid state/affects the cold crystallisation behaviour

favourably and might be exploited.

111

2. CHARACTERISATION OF RAW MATERIALS

The purpose of this section is to investigate the properties of

the raw materials used in the study. The results determined procedures

used and the compounds produced in the processing stage.

The effect of moisture was investigated along with drying procedures.

The effect of different nucleants was also studied in some detail.

2.1. THE RECYCLE BOTTLE

The bottles used in this investigation were supplied by Carters

Packaging Limited of Long Eaton. They were of polymer produced by

Eastman Kodak and are the type used for carbonated soft drinks (Kodapak

7352). Bottles of this type are not coated with PV-dC as they are not

used for sensitive products, and the shelf-life is sufficient. The

type of bottle used is shown in Fig 38.

The PET bottle derives its strength from the high degree of biaxial

molecular orientation that is composed during injection blow-moulding.

An example of its strength (97) is that a filled 21 bottle at a working

pressure of 4 atmospheres (typical for a carbonated soft drink) will

hold up to 12 atmospheres and still bounce when dropped from a height

of 3m on to concrete. This high strength gives several benefits to the

producer and consumer.

112

Fig 38: The two litre carbonated soft drinks bottle

PET

335 mm

, )

r 50 mm _____ BASE CUP

----- HIGH DENSITY POLYETHYLENE

1 ADHESIVE

r ~--------l00mm~·------~i-

113

iii

The typical properties of the bottle type used in this invest­

igation are shown in Table 18. Due to the high strength of biaxially

orientated PET very light weight containers can be produced.

PET is used pure in container manufacture - no additives are

required; this was confirmed by cast film spectroscopy. X-ray analysis

indicated that no heavy metal polymerisation catalyst residues were

present.

Extensive extraction testing. (97) in PET using a wide range of

food-simulating solvents has shown that PET poses no toxicological

hazard or other adverse effects to the human body. Indeed, PET is so

inert it can be used for various surgical applications.

The most common bottle shape used for carbonated soft drinks has

a hemispherical base design and requires a base cup for support. In

this case high density polyethylene is used as the support and to

absorb impact blows. The base cup is usually applied directly·after

stretch blow moulding of the bottle. The adhesive used was supplied

by Midland Thermoplastics Limited; adhesive reference 'HYTAK 43'. It

is stable to both temperature variation and solvents.

114

Table 18: Typical Properties of Oriented Sidewall Moulded

Property, Units

Wall Thickness,mm

Density, gcm-

Crystallinity, %

Intrinsic Viscosity, dl/9

Tensile strength at yield

Hoop, MPa

Axial, MPa

Tensile strength at break

Hoop, MPa

Axial, MPa

® KODAPAK PET (98)

Test Method

ASTM D 1505

ASTM D 881

Tensile modulus of elasticity

Hoop, GPa

Axial, GPa

Gloss at 45° ASTM D 2457

Water vapour transmission rate,

g/m'/24h ASTM E 96-E

Gas transmission rate

cm'/m'/24h ASTM D 1434

CO,

0,

115

Value

0.31

1.36

25

0.71 to 0.75

172.4

69.0

J93·1

117: 2..

4.275

2.206

100

2.3

3.1

6.2

2.1.1. HYDROLYTIC DEGRADATION OF PET: EFFECT OF RELATIVE HUMIDITY

ON THE MEASUREMENT OF THE MELT FLOW INDEX OF SCRAP PET

2.1.1.1. INTRODUCTION

It is well known and widely accepted that the principal cause of

loss of molecular weight in PET during processing is hydrolytic deg-

radation. By maintaining samples of regrind at various humidities it

is possible to see how moisture affects the molecular weight and thus

the melt flow index of the sample.

2.1.1.2. EXPERIMENTAL

The PET bottles were granulated to produce regrind, which was then

placed in dessicators with saturated solutions of salts to give different

relative humidities. (Table 19)

Table 19: Salts used to produce constant relative humidities at 20°C

(Source (99) Handbook of Chemistry and Physics)

Relative Humidity %

100

76

52

20

o

116

Salt

pure H,O

Na(CH,COO)3~0

Na, C, ~ 2H z 0

K(CH,COO)

Silica gel

The samples were left for 14 days to reach equilibrium before

testing; PET granules were also ke~t in the same environments.

Melt Flow Index of Poly(ethylene terephthalate)

The MFI was obtained using BS 2182 Method 105C (100,101) and ASTM

D 1238 (102). The PET, contained in a vertical metal cylinder,is

extruded through aJel:.B die by a loaded piston (2.16kg) at a temperature

of 270 o C. The apparatus was thoroughly cleaned between each test.

The cylinder was charged with ~ 2.5g of test sample. During the

charging operation, which took less than one minute, the sample was

continually tamped down using the charging tool. When charging was

complete the unloaded piston was inserted into the top of the cylinder.

Five minutes after inserting' the piston, the temperature returned

to 270 0 C and the load was placed on the piston to extrude the PET

through the die. The loaded piston was allowed to descend under gravity,

and the rate of extrusion was measured by taking a 10 second sample of

extrudate at the die. All the material extruded up to the point at

which the lower reference mark and the top of the cylinder were dis­

carded, as were any subsequent cut-offs containing air bubbles. The

remaining samples (4-8) were weighed individually. All of these samples

were taken when the piston was between 50mm and 20mm from the upper end

of the die.ie between the times when the first and second marks on the

piston disappear into' the cylinder.

MFI was calculated as the mass of extrudates in grammes per ten

minutes.

117

2.1.1.3. RESULTS

Table 20: MFI's of test materials (g/10min)

Relative Humidity PET(Virgin) PET(Regrind) ( %)

100 20.6 34.5

76 20.5 31.7

52 19.5 28.8

20 18.9 28.6

0 6.8 12.3

AS SUPPLIED 20.0 29.4

Generally the MFI of the samples increases with the relative

humidity. The difference between the dry PET and the rest is marked,

showing very clearly the importance of drying. It can also be clearly

seen that the PET granules - as supplied require drying before use.

It can be seen that the regrind has a lower viscosity than the

virgin. The regrind is more susceptible to hydrolytic degradation than

the virgin this is probably due to the regrind being in the form of flakes

(large surface area) whereas the virgin :material is in pellet form.

Subsequent work on the Davenport shear rheometer showed a similar

decrease in viscosity with increasing relative humidity. The work also

showed that PET and bis-phenol-A-polycarbonate are pseudoplastic (shear-

thinning) ~Dwer law fluids. Lexan 66 polycarbonate to be used in this

118

investigation has a viscosity ten times greater than the PET.

2.1.2. MEASUREMENT OF THE INTRINSIC VISCOSITY OF THE MFI EXTRUDATE (103)

2.1.2.1. INTRODUCTION

The intrinsic viscosity [1) of a solution is related to the viscosity

average molecular weight of the polymer Mv by the Mark-Houwink equation:

[ 1.) -a = KM"

where K and a are constants for a given polymer, solvent and temperature.

[1) is related to measurable quantities by the Huggins equation,

which may be written,

or

I/. sp c

= ['1.) + k'[t)2C + k'['Ll'c 2 +

(In'Zr) " = [t) - k'['l.)2 C + k ['Zl'c 2 + -- -----c

where 1r = relative viscosity of solution = t/to

1. sp = specific viscosity of solution = Z r -

t and to are measured flow times of solution and solvent

-3 and c= concentration of polymer solution in kgm

ie [~) is the intercept at coo of plots'1sp/c or (In~r)/c

against c at low concentrations.

11 a

2.1.2.2. PROCEDURE

One gramme of extrudate was placed in 100ml of a 60:40 by weight

solvent of phenol and tetrachloroethane. The solution was left to

stand for 2 days, then mechanically shaken and heated (50 0 C) until the

PET dissolved. Solutions were made up at concentrations of 1, 0.75, 0.5

and 0.25%.

2.1.2.3. RESULTS

Table 21: Intrinsic Viscosity of MFI Extrudate kept at different

Relative Humidities

Relative Humidity IV of Extrudate % dL/3

0 .55

20 .48

52 .44,

76 .41

100 .38

AS SUPPLIED .46 ,

BOTTLE .75

The bottle regrind has an intrinsic viscosity of 0.75 after

extrusion which suggests that the effect of the process on the

molecular weight of the regrind is minimal. The effect of moisture in

reducing the, 1. V is clear. The regrind at 0% RH has'-not' been' di'ied'"

120

effectively in the dessicator as the drop in I.V is marked. The intrinsic

viscosity is a sensitive measure of the moisture content of the .PET

regrind prior to extrusion. As the relative humidity increases there is

a detectable drop in the 1. V of extrudate ..

2.1.3. COMPARISON OF MFI AND I.V RESULTS

To enable comparison of intrinsic viscosity and melt flow index, 0.

MFI values were converted into mel,t viscosity data. Fig 39 shows/'comp-

arison of melt viscosity and intrinsic viscosity. The relationship be-

tween mel t viscosity as measured in the MFI apparatus at 270°C and

intrinsic viscosity of the extrudate is shown in Fig 40.

121

Fig 39: Comparison of Intrinsic/Melt Viscosities

Intrinsic

Viscosity (dL/jJ

Melt Viscosity

(MPa :0)

x 10-4

x

0.6

0.5 1.0 lr

0.4 0:8

0.3 0.6

o

0.4

0.2

o

o

25 50

RELATIVE HUMIDITY (%)

122

---_.-

o

'"

75 100

Fig 40: Melt Viscosity as a function of IV

Melt Viscosity

(MPa s) 'IO~

3

2

o 0.1 0.2 0.3 0.4 0.5

INTRINSIC VISCOSITY (dL/j)

123

0.6 0.7 0.8

2.2. EFFECT OF DRYING TIME ON THE MFI OF SCRAP PET

2.2.1. DRYING PET REGRIND IN A DESSICATOR AT 23°C

The melt flow index of the regrind is a very good indication of

the moisture content in the sample. The more moisture present the

larger should be the MFI as the viscosity decreases due to hydrolytic

degradation of the PET.

To measure the drying time,PET regrind was placed in a silica

gel dessicato~ and whilst drying, portions of the regrind were removed

the time noted and MFI measured.

The results obtained are shown in Fig 41. The initial values for

the MFI would suggest that the sampling intervals were too close to­

gether and that the influx of "wet" air on opening the dessicator was

counter-acting the drying by the silica gel. If this were taken into

account the drying of the regrind has a linear relationship with log t.

It can be seen that the regrind when as dry as possible (maximum

drying capacity of silica gel) it still has an MFI of 12.28.

2.2.2. DRYING IN AN OVEN AT 120°C

120°C was selected as the test temperature as this value is most

frequently quoted in literature.

PET regrind was placed in the oven and the melt flow index measured

at different time intervals. The results are shown in Fig 42. It can be

clearly seen that optimum drying time for PET regrind is ~ 9! hours.

124

Fig 41: Drying curve of regrind :PET in a dessicator at 23°C

30

20

10

- 0.4 o 0.4 0.8 1.2 1.6 2.0 2.4 2.8

MFI (;>;/10"";")

32 Fig 42: Drying curve of regrind PET in an. oven at 120°C

28

24

20

12

8

4

o 4 8 12 16 20 28 32 36 40 44 48

After this time the MFI starts to increase, this is probably due to

thermal degradation. Less than this time hydrolytic degradation is

the cause of the high MFI.

2.2.3. COMPARISON OF DRYING METHODS

Drying in an oven is the only way that drying of the granules

can be achieved; it would appear placing the granules in a dessicator

. I merely removes the surface moisture.

Drying granules in an oven for 9~ hours would not be feasible

in industry·. The solution is to use a dehumidifying oven which dries

granules in 2 to 2~ hours at 150°C.

2.3. STUDY OF CRYSTALLISATION UNDER VARIOUS IMPOSED TEMPERATURE

PROGRAMMES AND INVESTIGATING THE EFFECT OF NUCLEATING AGENTS·

The crystallinity of PET has an effect upon physical properties;

3enerally the more crystalline the PET the better its physical

properties. It is therefore important to see how PET is affected by

different treatments.

Observations were made using the Du Pont thermal analyser fitted

with a differential scanning calorimetry (DSC) cell. The temperature

difference between the sample and the inert reference is plotted against 1Nl te,.,peroJvtt- or

temperature sOAthermal changes such as crystallisation tenperature, Tc)

may be observed.

By measur:rrig:; the area under the melting peak, on heatingjor crystallisation

peak on cooling)the heat of fusion

E = _---Cf1::.;Hc::m'---_

can be determined from the following:

60 A B f1qs

where f1H = Heat of fusion, Jig

M = sample mass, mg

A = peak area in cm 2

B = TIME BASE setting, min/cm

eg 100 Umin = lmin/cm

5DClmin = 2min/cm

f1qs = Y axis sensitivity setting in mV/cm (5mV/cm)

E = cell calibration coefficient in mW/mV; in

this case 0.22 mW/mV

The crys tallisation behaviour was examined by imposing the

following regime on the PET specimen

Sample size, 8-10mg in a Nitrogen atmosphere

Heat at 10 deg C/min to 270DC

Maintain at 270 DC for 5 min

Cool at 5 deg C/min to 30-40DC

Reheat at 10 deg C/min to 270DC

The variables covered were:

effect of processing

addition of 0.25 and 1% talc

addition of 1% hammer milled glass with surface treatment

addition of other nucleating agents at 0.25% level

addition of bis phenol-A-polycarbonate at 20% level

Once the heats of fusion and crystallisation had been obtained it

128

.was possible to calculate the percentage crystallinity of the sample.

The heat of fusion of hypothetical 100% crystallinepolyethylene

5 -1 terephthalate (6Hf 100(PET)) is 1 .22 x 10 J.kg. The results obtained

are shown in Table 22.

The effect of drying PET at 120°C did not increase the crystallinity

very much on heating, but on cooling the absence of water increased the

crystallinity from 20.29 to 29.30%. Crystallisation also occurred at a

higher temperature in the thoroughly dried sample. This would suggest

that water has an effect on crystallinity as well as causing hydrolytic

degradation.

Processing has very little effect upon crystallinity. Talc is

recommended in the literature as a nucleating agent for PET. Indeed

the initialcr-ystallisation and crystallisation temperatures are highest

of this nucleant however the crystallinity of this compound is not in-

creased.

Hammer-milled glass increases the crystallinity which is encouraging

as during processing some of the glass fibres used will be broken down

and if these contribute towards nucleating the PET this is an added bene-

fit.

® Pansil which has a magnesium silicate base (c.f talc) does produce

a high degree of crystallinity but the crystallisation temperatures are 0.5 ,$e.,.

notLhigh asktalc. In future a combination of Pansil/talc nucleant might

prove effective.

129

Table 22: PS! Crystallinity

COMPOUND ~ CRYSTALLINITY To % CRYSTALLINITY Ts Te heating ·c cooling ·C ·C

PST

PET dried 23°C 24.8'<. 244 20.29 185 165

PET dried 120°C 26.30 248 29.30 206 183

eXTRUDED PET

BETOL SINGLE SCR~~ 27.05 248 21.05 211 194

B-P TwIN SCRB-I (245'C) 33.28 248.5 29.96 209 197.5

9-P TWIN SCRB, ( 2100 Cl 28.02 247 28.68 207 195.5

PET/PC 80120 34.49 248 16.23 192 177

PS'I' /KM334 95/5 20.09 248 18.55 207 194

PSi' + Additives·

0.25% talc 28.53 247 29.72 218 208.5

1% talc 28.51 246 33.54 208 193.5

1% hammer-milled glass 31.38 248.5 33.54 208 193.5

PET nucleants

SiO t Insert substances (0.25%) 28.40 250 27.05 207 191

TiO~ 32.38 250 35.78 210 191

P.o.NSIL 43.98 251 31.41 213 199

ORGANIC COMPOUNDS (0.25%)

dimenthyl isophthalate 28.72 250.0 31.60 212 196

terephthalic acid 34.54 250.5 26.63 214 197.5

sodium acetate 41.16 251.5 28.41 217 209

sodium stl3rate 30.57 250 42.60 215 207

diphenyl amine 31.59 251.5 31. 59 210 194

POLY!-1ERS (0.25%)

polypropylene 34.49 248 16.23 215 199

LWP':: 33.94 248.5 31.82 209 190

-r;;.. # ~t:r~J ~~'re

7S # ~~ 'If ~,~ ~toJ,l.i...t<Mo., -re. 0 '4jSiaiJ.<o a.t; "'" "-yerrdwre

130

Of the numerous nucleants tried, the salts of c~r.boxylic acids

were the most promising. Sodium acetate and stearate increased

crystallisation greatly in heating and cooling respectively. Sodium

stearate was most useful as a nucleant as it promotes crystallisation

in specimens whilst cooling in a mould. Sodium stearate starts

crystallising at a high temperature (Ts = 215°C), has a high cryst­

allisation temperature (Tc = 207°C) and has a fast rate of crystal­

lisation. These factors are all useful in injection moulding PET for

decreasing the cycle time.

1 31

3. PROCESSING.

3.1. COMPOUNDING

3.1.1. INTRODUCTION

Compounding was performed on a Baker-Perkins MPC/V 30,

the laboratory model twin screw extruder. Technical specifications of

this model are given in Table 23.

Table 23: Technical Specification for Baker-Perkins MP C Iv 30

Unit

Length: diameter ratio 13,1

Drive power, kW 3

Maximum screw speed (rpm) 500

Barrel diameter (mm) 30

Number temperature control zones 3

Ins talled barrel heating (kl,) 5

Barrel cooling (Kcal/h) 5040

All resins were thoroughly dried.

132

3.1.2. SCREW CONFIGURATION

In this case the extruder was being used as a two stage mixer.

In the first stage PET was combined with polymer (PC) or impact

modifier (KM 334) and in the second glass fibres were added. This

requires two mixing sections below the feed barrel orifices (FBO),

these mixing sections were the main consideration in designing the ;

screw. These mixing sections were made up of 600~ feed paddles (FP)

which have some forward conveying tendency. Orifice plugs (OP) are

used to hold melt in a mixing section. Feed screws (FS) were placed

to convey the materials between mixing sections. The feed screw tends

to push material through paddle sections(which have low forwarding

capabilit0, and over orifice plugs.

Feed screw spacers (FSS) were used where a conveying screw is desired,

but spacing of components does not allow for a standard one diameter (lD)

screw. The barrel valve is always used in conjunction with a pair of

orifice plugs directly beneath the barrel valve vane.

The screw configuration inserted was:

D~ FSS, D1 FS, D1~ FS, 8 x 60 0 FP,

(15mm x 29;85, FBO) D~ FSS, D1~ FS, (15mm x 29.44 FBO/BV), 5~~if f~

D1 FS, 1 x 00 P, D1~fS]00P, CB (Connecting Bolt)

133

3.1.3. OPERATING PROCEDURE

The twin-screw extruder was used as a two stage mixer with

volumetric screw feeders metering the feed into the ports. The

polymer was mel ted in first section and ran starved under the second

feed port half way down the barrel where the glass fibres were added.

The polymer is m~lted in the first stage and the fibres gently blended

in the second stage to minimize breakage.

Operating Conditions:

Screw Speed 200 rpm

Torque 70%

Barrel Control Melt Temperatures

256 255

262 261

259 265

265 DIE 270 DIE

Output 10-12 kg/hr

NB The extruder must be running before the feeders are started.

134

3.1.4. COMPOUNDS

Table 24: COMPOUNDS PRODUCED

Compound Polymer Reinforcemen t . Impact Modifier

PET(regrind)

2 PET(regrind) c 30 w/w % glass fibres

3 PET(virgin)

4 PET(virgin) 30 w/w % glass fibres

5 PET(virgin) 50 w/w % glass fibres

6 PET(virgin)+ 30 w/w % glass fibres sodium stearate

7 PET(virgin) 30 w/w % hammer-milled glass + silane treatment

8 PET(virgin) 30 w/w % hammer-milled glass

10 PET/PC 80120

11 PET/PC 80120 30 w/w % glass fibres

.. 12 PET(virgin) 5% KM334

13 PET(virgin) 10% KM334

14 PET(virgin) 30 w/w % glass fibres 5% KM334

These compositions were all produced without difficulty.

/I{~re£I -Mdo.MoJ{l ~ R6h,~ /l.f\cI It«c..S

- KM HIf lA ~ S\A.(.(UAW J::" KM 330 ·wh,;..iA wC<4 V«..d ad

1 3~

3.2. INJECTION MOULDING

The properties of a short fibre reinforced thermoplastic are very

dependent on fibre length and orientation it is important that both of

these parameters can be controlled in the final moulding, by an approp-

riate choice of moulding conditions (Table 25).

Table 25: Moulding Conditions for Bipel Injection Moulder

Melt temperature (settings): 250, 255, 260°C (barrel)

270°C (nozzle)

Mould temperature 140°C

Injection pressure 500 psi (3·45 MN';;2.)

Injection/Hold-on time 60s

Cooling time 290s

Screw rotation speed 130 rpm

The polymer used was dried and maintained at low moisture levels

during processing by using a hopper drier. Bars 190 mm x 12.5 mm x 3mm

were moulded without difficulty for all compounds.

136

4. TEST METHODS

4.1. FLEXURAL TESTING

It was decided to test the compounds in flexure in accordance

wit~ ASTM D790 (104) (for example as an automobile tail-gate) this

is the type of stress to which the component would be subjected. A

standard three point loading system, centre loading on a simply

supported beam, was used.

The Instron was used with a 100kg load cell (tension). After

calibrating the load cell the flexural tests used a crosshead speed of

10mm/min7· The chart speed was set at lcm per 2mm of crosshead. Five

speCimens were tested for each compound.

The depth of the beam tested was 3.15mm and the span 60mm

(L/D=19), this follOWS ASTM D790 as the support span should be 16

(tolerance +4 or -2) times the depth. The speCimen was long enough

(190 mm) to allow for overhang at each end greater then 10 percent of

span.

The advantage of the Instron was that the flexural rig fits inside

the oven, so.specimens could be tested at diferent temperatures (23, 40,

and 150°C). The specimens were conditioned for 56 hours at 23°C at 50%

relative humidity prior to testing, and tested under the same conditions 1

except for heated samples which were kept at temperature for 20 minutes

before testing.

\

137

Equations used were:

Maximum flexural stress,

S = 3PL

2bd'

Modulus of elasticity,

E = L'M B 4bd'

Maximum fibre strain,

r = 6Dd V-

where

P = maximum load of, N

L = support span, m

b = width of beam tested,

d = depth of beam tested,

M = slope of the tangent

hi

hi

to the

S = stress in outer fibres at

midspan, NI"'"

EB = modulus of elasticity in

bending, N/m'

r = maximum strain in the outer

fibres, mm/mm

(D,d and L in mm)

initial straight-line portion of

the load deflection curve, N/m

D = maximum deflection of the centre of the beam, mm

4.2. IMPACT TESTING

The impact strength of a material to be used for automotive purposes is

important. The relevance of this test was to determine the resistance of

the compounds to flexural shock. The Ceast "Advanced Fractoscope" was

used to test the specimens, with a Charpy type test, with the specimen

supported as a horizontal simple beam and broken by a single swing of the

pendulum.

The pendulum contains a transducer which enables the force/time character­

istic to be derived.

The tests were performed on unnotched samples as the testing was

merely to compare impact strengths of the compounds, and not to study

fracture mechanics which is beyond the scope of this programme.

The cross-section of the test specimens was 'V 3.15 x 12.50mm.

All specimens were conditioned at 23°C and 50% relative humidity before

being tested at -20 and 40°C. Five specimens were used for each compound

to be tested at each temperature. The span was 48mm; L/D~ 15.24.

The impact test was conducted at a constant test velocity of

Vo = 3.46 m/s the mass of the tup was 4.17 kg and therefore the impact

energy available was 24.96J. The Ceast Impact Tester is instrumeted and

there was no necessity to carry out further calculations. The information

obtained from the instrument was Max Force (N)

Energy to Failure (J)

and Total Energy (J)

139

4.3. ENVIRONMENTAL STRESS CRACKING

The specimens to be tested were placed in a flexural rig at the

following constant outerfibre strains (E) 7;31. 5.08, 3.30, 1.90 and

0.25%, and subjected to the cracking agents at the midspan for 24 hours.

The agents were selected particularly for their relevance to the

automotive industry: petrol, salt solution and water (as a control).

After 24 hours the specimens were examined for cracks or crazing.

4.4. ANALYSIS OF FRACTURE SURFACES

The specimens to be examined were gold splutter-coated after drying

at 120°C for 3 hours. The fracture surface was examined for mode of

failure and any orientation effects using the scanning electron micro­

scope. (Cambridge Stereoscan Mk2).

4.5. FIBRE LENGTH

Fibre length was examined to discover the amount of fibre breakage

that had occurred during processing. This was determined by first burn­

ing off the PET polymer in a furnace at 600°C and then examining the

residue under a microscope to determine the average fibre length and

fibre length distribution.

140

5. EXPERIMENTAL··RESULTS AND DISCUSSION

5.1. FLEXURAL PROPERTIES

5.1 . 1 . MAXIMUM FLEXURAL STRESS.

Compound Polymer Reinforcement Maximum Flexural Stress (MPa) (Standard Deviation)

150°C

PET (regrind) 106.26- (5.831 80.51- (5. 71)

2 PET (regrind) T30 w/w% glass fibres , 13.26 (4.40) '126.90 (2.771

3 PET (virgin) 94.81 ( 1.68) 84.27 (1. 78) 14.22 (D.4D)

4 PET (virgin) T30 w/w% glass fibres 121.49 (5.29 123.03 (3.26) 55.68 (D.61)

5 PET (virgin) T50 w/w"/. glass fibres 110.26 (5. 40) 1 15.45 (5.93) 80.98 !4.101

6 PET (virgin) • +30 w/Wfo glass fibres (0.25%) 111.22 (3.46) 116.15 (J.50) sodium stearate

7 PET (vi.rgin) 30 w/~ ha'ttlller-milled glass T silane treatment 90.44 (2.66) 88.23 (2.14 )

8 PET (virgin) )0 w/w"J. hamner-milled glass 84.78 (4.25) 83.22 (2.05)

10 PETIPC 80/20 83.84 ( 1.53) 84.21 (3.351 5.84 (0.35)

11 PETIPC 80120 30 w/w% glass fibres 141.61 (6.71 ) 143.24 (4.42) ·45.19 (2.13)

12 PETIKM 334 95/5 14.43 (0. 58} 66.24 (2.D2) 8.67 (0.20)

13 PET/KM .334 90/10 51.09- (0.83) 52.85 (0.05) 8.61 (O.H)

14 PET/KM334 95/5 30 w/w% glass fibres 111.09 ' (5.22) 108.76 (5.16)

-Corrected for large deflection

These values are considered correct for AIL =0.1 and Lld =16. The valve of AIL is the limit of

deformation which can be reasonably· tolerated without introducing corrections. As Lld = 19 in this

case there should be a correction made of +2.2% to the values of fleJCUral stress.

1 lil

L ~ Sp3n 60rrn 5.1.2. DEruI::!"Ial rJ: EIlF.OK

Om.s:TICN (m) M. a>1RlID Rl!..lMEll 23°C 4O"C 15O"C 23°C 4O"C 15COC

PCr (nw'.n:l) 13.9 9:' 0.232 0.152

2 PCr (nw'_'1d) 30 wM glass fibres 3.93 4.8 0.C>i6 O.tro

3 PCr (virgin) - DID rm 8f!EAK

4 PCr (virgin) 3D wN'l. glass fibres 4.6 4.7 '" ffiFAK

O.cm 0.078

5 PCr (virgin) 50 w/ ... <>f,. glass fibres 2.7 2.5 4.6 0.045 0.042 O·cm

6 PEl' (virgini + 30 w/t.."Io glass fibres 4.8 5.7 0.000 0.095 s:x11un stearate

7 PCr (virgin) 30 wlw'lo ~!"'~ed 6.5 6.6 0.103 0.1'0 glass + silare treatm;nt

8 PCr (virgin) 30 w/W'/. ha1:trer-iDil.led glass 6.5 6.5 0.1(77 O. lOO

10 PEI'/FCOO/20 OID tor EIlF.OK

11 PErt?: 00120 30 '.If.,.'''/.. glass f1bres 6.3 6.6 m BREA.I( 0.105 0.110

12 PCrIKMJ34 g;15 DID rrn: BREAK

13 FEI'1RM334 9:)/10 8.5 6.7 0.142 0.112

14 PCrIKMJ34 g;15 JO w/vi'/. gl2ss fibres 5.8 6.2 o·rm 0.103

142

5.1.3. FLEXURAL MODULUS

Flexural Modulus,GPa (Standard Deviation)

Compound Polymer Reinforcement 23°C 40°C 150°C

PET ( regrind) 2.10 (0.01) 1.80 (0.031

2 PET (regrind) 30 w/w10 glass fibres 6.04 (0.201 5.52 (0.161

3 PET ' (virgin) 2.00 (0.091 2.04 (0.101 0.40

4 PET (virgin) 30 w/w1. glass fibres 5.31 (0.191 5.15 (0.251 1.48 (0.061

5 PET (virgin) 50 w/w% glass fibres 9.63 (0.051 9.63 (0.361 5.01 (0.081

6 PET (virgin) • 30 w/w"lo glass fibres 4.50 (0.201 3.91 (0.06 sodim stearate

7 PET (virgin) 30 w/w% hamme~-milled glass ~ 3.36 (0.081 2.98 (0.14) silane treatment

8 P<--, (virgin) 30 w/Y% ha~er-milled glass 3.27 (0.151 3.00 (0.07)

10 PET/PC 80120 1.81 (0.01 I 1. 79 (0.051 0.14

11 PST/PC 80/20 30 wN% glass fibres 5.02 (0.251 5.07 (0.251 1.31 (0.021

12 PET 11<1-L334 95/5 1.66 (0.081 1.61 (0.051 0.30 (0.01 I

13 PET/!a: '334 90/10 1.51 (0.071 1. 49 (0.031 0.43 (0.02)

14 PET IKM334 95/5 30 w/w% glass fibres 4.00 (0.11 ) 3.93 (0.081

143

5.1.4. Fibre Strain

Compound Polymer Reinforce:nent Maximum ~ibre Strain lmm/mmJ

PeT ( regrind) 2]OC 40°C 150°C .0773 .0523

2 PET ( regrind) ]0 w/w% glass fibres .0215 .0255

3 PeT (virgin) .0700* .0981* .15]8*

• PeT ('/irgin) ]0 w/ift. glass fibres .D2£.:- .0250 .0£.79*

5 PeT (virgin) 50 w/w"/. glass fibres .0'149 .0137 .0253

6 PET (virgin)+ 30 .. lift. glass fibres .0256 .0306 sodium stearate

7 PeT (virgin) 30 w/w"J. ham:ner-milled .0346 .0352 glass + silane treatment

8 PET (virgin) 30 w/wo/. hammer-milled glass .0337 .0343

10 PETIPC 80/20 .0458 lt .0455* • 11 PETIPC 80/20 30 w/ift. glass fibres .03]7 .0359 .0947*

12 PETIKM.:334 95/5 .0510* .0678* .0'(50*

13 PET/KM:": 334 90/10 .0457 .0351 .0770

" PET IKM334 95/5 30 w/ift. glass fibres .0304 .0326

fOid not break

PET regrind has a similar (maximum) flexural strength to virgin,

but is inferior in strain accomodation as the regrind specimens fracture {" -

whilst the virgin specimens ,:cIo <loot.

The regrind with 30 w/w% glass fibres compares extremely well with

virgin + 30 w/w% glass fibres. Hammer-milled glass is inferior to glass

fibres in PET composites, hammer-milled glass has 60% of the stiffness of

the composite with glass fibres. The properties of the hammer-milled

glass are not improved by surface treatment.

The PET/PC 80/20 blend is similar in flexure to PET at 23°C and 40°C

but inferior at 150°C. The PET/PC with 30 w/w% glass fibres had the

greatest flexural strength of all the compounds at 23°C and 40°C, but was

inferior to PET reinforced with 30 and 50 w/w% glass fibres at 150°C.

The impact modifier KM334 reduced the flexural stiffness considerably.

145

5.2. IMPACl' fRlPERI'IES

~ Ebl.,.,er Reinforc::erent M3ldr.u:l Force ~ .J.

Erergy to ist Failu.'"'e 'lbtal Ere."'gf

l<tIiC s"""""' J J r:evtat1cn)

-2O'C 4CI'C -2O'C WC -2O'C 4U'C

PErlregMn:11 86 131 85 13.61 .6C .77 .69 .76

2 PErI """,.rd I 30 wM glass fibl"'eS 131 161 137 13.81 .82 .77 .86 .72

3 PErlvirglnl 12316.61 124 (7.31 1.68 1.69 1.69 1.68

4 PErI virgin I 30 wN'/. glass fibres 13416.21 133 (7.21 .77 .82 .81 .81

5 PEr I virgin I 50 w/...tj. glass fibres SEE FIG 43

6 PEr{ virgin) + 30 wM glass fibres 12313.71 141 (6.8) .82 .!J) .82 .!J) s::diun stearate

7 PErlvirglnl 30 wN/. h:mtEr-m:illed 11315.71 111 (6.0) .77 .70 .75 .69 glass + si.lare treahlalt

8 PErlvirglnl 30 wr..lo hame!'-milled 116(6.0) 10') 12.41 .82 .71 .83 .67 sl=

10 PEr/!C W20 141(7.3) 14317.21 2.31 3.0') 2.35 3.1J7

11 PEr/!c .00120 30 '01/ ... " glass fibres 17218.71 1!J) 18.351 1.67 1.10 1.11 1.70

12 PEr/OO34 9515 9214.31 93 12.81 .84 .82 .i57 .!J)

13 PEI'M1334 90/10 75D.ll 7212.91 .54 .47 .58 .')()

14 PEr/0034 9515 30 w/YIfo glass fibres 11917.01 11516.8) .76 .79 .82 .85

.. E cU-~ 1""- -h.-.,A. ~ ~du.te.

+ Totr..l &WO} tk I1Nl 'i 1Nl. t"W"\

I4-\:>

Fig 43. Impact Strength dependence on moulding thickness for 50 w/w% glass fibre

reinforced PET

Force (dN)

20

100+---------------~--------------~--------------~----------------~--------------~ 3.10 3.15 3.20 3.25 3.30 3.35

Thickness of Moulding (mm)

Regrind PET is inferior to virgin in impact, having 70% of the

maximum force, and 40% of the energy. These differences are no

longer.evident in composites with 30w/w% glass fibres which are similar.

Hammer-milled glass is inferior to glass fibres in strength of

PET composites. Hammer-milled glass shows no improvement over unrein­

forced whereas there is a 30% increase for glass fibre reinforcement at

the same level. The performance of hammer-milled glass is not improved

by surface treatment.

The PET/PC blend gave the highest results for impact (maximum force

and energy to first and total failure) of the compounds covered except

for the PET/PC blend with 30w/w% glass fibres which gave a higher

maximum force. The impact energy of the blend was 50% better than the

next best.

The impact modifier KM334 did not improve the impact properties of

PET.

The impact properties of PET + 50w/w% glass fibres were dependent

on the thickness of the mOUlding. The thicker mouldings had reduced

impact strength. It is usual for stiffer specimens as they get thicker

to tend towards brittleness. This would suggest that thicker mOUldings had

;. '.> 50% loading and fibre/fibre interaction was the mode of failure.

148

5.3. ENVIRONMENTAL STRESS CRACKING

CONSTANT STRAINS % 7.31, 5.08, 3.30, 1.90, 0.25

PETROL

SALT SOLN.

WATER

PET(VIRGIN)

NONE

NONE

NONE

* Only 0.25% strain used.

PET(VIRGIN) + 30 w/w%

GLASS FIBRES

NONE*

NONE*

NONE*

PET/PC

80120

NONE

NONE

NONE

PET/KM334

90/10

NONE

NONE

NONE

None of the compounds tested exhibited stress cracking. This is

particularly significant for the PET/PC blend as bis phenol-A-polycarbonate

is known to be particularly sensitive to environmental stress cracking by

petrol. Blending with PET however appears to protect the PC and the blend

is not weakened when placed in the hostile environment.

149

5 . 4 . EXAMINATION OF FRACTURE SURFACE USING THE SCANNING ELECTRON MICROSCOPE

Note: all surfaces examined were transverse to flow

5 . 4 . 1 . PET + 30 w/w% GLASS FIBRES

Fig 44 . Scanning electron micrograph showing fibre pull- out : vertical

glass fibres and pits .

Fig 45. The tip of fibre showing polymer attached that came away

during pull-out

Fig 46. Fibres bound together in a bundle by the polymer

151

Fig 47 . * A glass fibre showing the surface treatment . Note the pits

produced by fibre pull- out and the polymer that has become

attached to the fibre during the " explosive" fracture .

152

5 . 4.2 . PET+50 w/w% GLASS FIBRES

Fig . 48 The fracture surface showing vertical fibres at the edges

and more horizontal fibres in the centre (perpendicular to flow)

1~1

Fig . 49 Fibres at the edges mostly vertical , parallel to the

direction of flow.

Fig . 50 A dense clump of fibres showing clearly the "pull-out" mode

of failure .

154

Fig 51. The tubes left after fibre pull- out

155

For PET +30 w/w% glass fibres the fibres all appear to oriented

vertical on the fracture surface which suggests that the lie in the

direction of flow. The orientation of the fibres for 50% glass fibres

is more complex than in the 30 w/w% composite. The overall appearance

is more random; at the edge the fibres are vertical to the fracture

surface, but at the centre they are generally horizontal. This would

suggest that in flexure most of the stiffness is provided by the fibres

at the edges.

The mode of fracture for both the fibre loadings appears to be fibre

pull-out with very little breakage. This means that the full potential

of the system has not yet been realised. The reasons for this are; the

fibre length (discussed later) is too short to prevent fibre pull-out • 11:

and the surface treatment on the glass fibres (seen in the SEM's) is not

bonded sufficiently strongly to the polymer. On fibre pull-out the break

takes place between the surface treatment and the polymer.

It should be remembered that during impact fibre pull-out makes a

large contribution towards impact strength.

l~ii

--- -------

5.5. FIBRE-LENGTH-DISTRIBUTION

The glass fibre length distribution is shown in Fig 52, the mean

fibre length is 0.29 mm. The fibre length has a considerable effect

upon the physical properties of the composite.

The strength of a glass fibre composite is increased by long fibres

ie those greater than critical fibre length lc. (see p 11 ). This is in

contrast to composite toughness1which is increased by fibre lengths < lc.

This anomally can be settled by having a fibre lengths less than and

greater than the critical length.

The result of this would be that fibres shorter than lc would be

pulled from the matrix, when a crack passes through the composite,

energy absorbed and toughness increased. The long fibres would contribute

towards stiffness and strength provid~- the interfacial bond quality was

sufficient.

157

Fig 52. Glass fibre length distribution obtained using

a semi-automatic image analysis procedure

Frequency

60

50

40

30

20

10

o .05 . 1 .15 .2 .25 .3 .35 .4 .45 .5 .55 .6

fibre length (mm)

158

'\ /

)

6. CONCLUSIONS

1. The results of this investigation have shown that regrind

polyethylene terephthalate could be the basis of a reinforced

blend/composite that may be of use to the automotive industry.

Although the regrind is somewhat brittle, when 30 w/w% glass

fibres are added it behaves no differently from the virgin.

2. When PET is compounded with· 20% bisphenol-A-polycarbonate

a blend is produced which exhibits similar flexural properties

to the unmodified PET. The results for impact testing of this

blend show that this toughened blend should be suitable for a

bumper application. This blend does not suffer from environ­

mental stress cracking in petrol.

3. Reinforcement with 30 w/;t~ glass fibres of the 80/20 PET/PC

blend produces a composite exhibiting higher impact strength than

the equivalent (untoughened) PET plus 30 w/w% glass fibres. The

features of the composite/blend could be exploited in producing

body panels for automobiles, eg tail-gates.

SUGGESTIONS FOR FURTHER WORK

1. In view of the increase in properties of PET containing up to

50% glass fibres, it is recommended to repeat the experiments with

80/20 PET/PC in order to achieve increases 'in high temperature and

impact properties.

2. Close examination of the fibre-matrix interface should be under­

taken with a view to establishing the bond strength and whether a

more suitable coupling agent for glass fibre reinforced, PET/PC

blends is required.

160

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