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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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