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Original Article
Preparation, microstructures, and
properties of long-glass-fiber-reinforced
thermoplastic composites based on
polycarbonate/poly(butylene
terephthalate) alloys
Yahui Tan, Xiaodong Wang and Dezhen Wu
Abstract
This article investigated the preparation, microstructures, and performance of long-fiber-reinforced thermoplastic
polycarbonate/poly(butylene terephthalate) alloys through the joint processing of melt blending followed with a melt
pultrusion. The mechanical evaluation demonstrated that the resulting composites achieved significant improvements in
tensile strength, flexural strength and modulus, and notched impact strength. Such a prominent reinforcement effect is
attributed to the feature that the residual fiber length within the injection-molded long-fiber-reinforced thermoplastic
composite specimens is much longer than that of the short-fiber-reinforced ones. This takes full advantage of the
strength of the reinforcing fibers themselves. The improvement in impact toughness is due to the energy dissipation
by both the fiber pullout and fiber fracture as a result of the long-fiber-reinforcing effectiveness. The scanning electron
microscopy investigation confirmed that the fiber pullout and fiber breakage concurred on the impact and tensile fracture
surfaces, and the former was more significant than the latter. Meanwhile, the scanning electron microscopy observation
also indicated a good interfacial adhesion between the thermoplastic matrix and fibers, which results in the subsidiary
enhancement of mechanical properties. The study of dynamic mechanical analysis revealed the long-fiber-reinforced
thermoplastic polycarbonate/poly(butylene terephthalate) composites achieved a remarkable increase in storage modu-
lus but presented a considerable decrease in loss-factor magnitude compared to the pristine alloys. The thermal
stabilities of the composites were also slightly improved in the presence of long glass fibers according to the results
of thermogravimetric analysis.
Keywords
Long fiber reinforcement, polycarbonate/poly(butylene terephthalate) alloys, composites, mechanical properties,
morphology
Introduction
Fiber-reinforced thermoplastic composites are a class
of the most important engineering materials which
mainly benefit from the combination of a polymer
matrix and inorganic fiber. They cannot only provide
stiffness and impact resistance at a very light weight but
also give additional weight and/or cost savings for
equal or better performance compared with metal
materials.
1,2
The short-fiber-reinforced thermoplastic
composites have been well developed in the past cen-
tury, and their industrial applications have also been
rapidly growing due to their remarkable mechanical
performance as well as some interesting features such
as the capability to be molded in complex geometries,
State Key Laboratory of Organic–Inorganic Composites, Beijing
University of Chemical Technology, Beijing, China
Corresponding authors:
Xiaodong Wang, State Key Laboratory of Organic–Inorganic Composites,
Beijing University of Chemical Technology, Beijing 100029, China.
Email: wangxdfox@aliyun.com
Dezhen Wu, State Key Laboratory of Organic–Inorganic Composites,
Beijing University of Chemical Technology, Beijing 100029, China.
Email: wdz@mail.buct.edu.cn
Journal of Reinforced Plastics
and Composites
2015, Vol. 34(21) 1804–1820
! The Author(s) 2015
Reprints and permissions:
sagepub.co.uk/journalsPermissions.nav
DOI: 10.1177/0731684415599071
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low manufacturing cost, high production rate, and sig-
nificantly low weight-to-strength ratio.
3,4
However, the
structural applications of short-fiber-reinforced mater-
ials tend to be limited when compared with the com-
posites reinforced with continuous fibers, because their
mechanical performance is remarkably affected by the
size and orientation of the fibers in the final product.
5
In addition, the shear forces inherent in the process
promote a significant reduction in length of the fibers
during the injection molding of short-fiber-reinforced
composites, leading to a decrease in final strength. In
this case, the research and development of the new
reinforcement and processing technologies has
attracted tremendous attention over the past decades.
The long-fiber-reinforced thermoplastic (LFT) technol-
ogy is considered as an attractive alternative for creat-
ing a wide variety of high performance composite
materials, because it can bridge the gap between
short- and continuous-fiber thermoplastic composites,
offering better mechanical properties than short-fiber
materials but retaining the ability to be injection
molded.
6
Apart from their excellent economic benefits
and high productivity in processing, the LFT compos-
ites exhibit unique properties in mechanical strength
and modulus, impact resistance, fatigue resistance,
and durability, which are superior to the short-fiber-
reinforced ones. Nowadays, the LFT composites have
become well established as high performance engineer-
ing materials for structural applications in aviation,
aerospace, automobile, wind electricity, yacht, civil
engineering, sports and leisure products, medical equip-
ment, and major industrial areas.
7–9
The basic processing technologies for LFT compos-
ites have also been described quite well in the literature,
and the LFT composites are usually manufactured
through melt pultrusion, where continuous filaments
of glass fiber rovings pass through a thermoplastic
resin impregnation unit. The LFT composites can gen-
erally be classified as the LFT direct products and LFT
granular products.
10,11
The former are obtained
through online or direct compounding of LFT compos-
ites, in which the processor compounds the glass fibers,
polymers, and other additives in line with the molding
or part extrusion process. The latter are just the semi-
finished LFT materials in granule or pellet form which
are provided as bundles of fibers preimpregnated with a
matrix polymer.
12
With combining the advantages of
the direct LFT products and short fiber-reinforced
thermoplastics, the LFT granular products represent
the state of the art and are suitable for both the classic
injection-molding process and the injection compres-
sion molding process to fabricate parts with compli-
cated shape and structure. Such a type of LFT
composites have recently received much attention and
are finding ever-growing applications due to their
excellent short- and long-term mechanical perform-
ances compared to their challengers.
13
Although other
long fibers such as carbon, aramids, and natural fibers
are used in specialty applications, long glass fibers are
commonly employed as the reinforcing materials for
the fabrication of the LFT granulating composites.
14,15
The glass fibers are completely enveloped by this mater-
ial and then subsequently cut into pellets of a certain
desired length around 10–25 mm.
16,17
Such a processing
technology has also been described quite well in the
literature.
18–20
On the other hand, the thermoplastic
polymers for the process of LFT granulating compos-
ites are most commonly limited to polypropylene (PP)
and polyamide 6 and 66 due to their good fluidity and
processability.
21–25
Other applicable thermoplastics
include polyethylene terephthalate (PET), poly(buty-
lene terephthalate) (PBT), acrylonitrile–butadiene–styr-
ene copolymer (ABS), high-density polyethylene,
polyphenylene sulfide, and thermoplastic polyureth-
ane.
11,26
Globally, about 65% of the LFT composite
market is PP based, and polyamide 6 and 66 have
about 20% of market share, whereas other resins
such as PET, ABS, and high temperature polyamides
comprise the remaining 15%.
27,28
Polymeric blends or polymer alloys, by definition,
are physical mixtures of structurally different homopo-
lymers or copolymers. In this case, the mixing of two or
more polymers provides a new material with a modified
array of properties. The blends of thermoplastic poly-
esters, namely PBT, with polycarbonate (PC), consti-
tute an important commercial polymer alloy, in which
the semicrystalline PBT provides chemical resistance
and thermal stability while the amorphous PC gives
impact resistance, toughness, and dimensional stability
at elevated temperatures.
29,30
Many studies indicated
that the PC/PBT alloy reinforced with short glass
fibers to enhance their mechanical properties.
31,32
However, it was also observed that considerable fiber
damage occurred during the injection-molding process,
thus reducing the average fiber length and decreasing
the reinforcement efficiency.
33–35
Moreover, the fiber
orientation and length distribution, interfacial bond
strength, and glass fiber concentration were also
responsible for the variation in mechanical properties
of the resulting composites.
23,36
On the other hand, the
mechanical performance of LFT composites is directly
connected to the relative fiber length of the reinforcing
fibers used. Although the original length of glass fibers
in the LFT granulating composites was usually over a
value of 10 corresponding to an effective fiber length of
1.2–1.4 mm, there was little consistency with regard to
the residual fiber length in the finished part, because the
residual length was finally determined by the molding
methods.
37,38
In this work, we prepared the LFT gran-
ulating composites based on PC/PBT alloys through
Tan et al. 1805
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the melt pultrusion of glass fibers. It is expected that the
LFT technologies performing upon the PC/PBT alloys
can lead to a much better reinforcement effect in com-
parison with the short fiber reinforcement and thus can
greatly exploit the applications of the PC/PBT-alloy-
based composites. The aim of this study is to clarify
the influence of the residual fiber length and fiber con-
tent on the mechanic properties and microstructures of
the resulting composites and thus to develop a class of
new high-performance LFT granulating composites
based on the PC/PBT alloys.
Experimental
Materials
The PC resin (Panlite
Õ
L-1250Y) with a specific gravity
of 1.2 and a melt flow index of 9.6 g/10 min was pur-
chased from Teijin Chemicals, Tokyo, Japan. The PBT
resin (Valox
Õ
315) with a specific gravity of 1.31 and a
water absorption rate of 0.08% (24 h, 23
C) was com-
mercially supplied by Sabic Innovation Plastics
Holding BV, Saudi Arabia. A solid epoxy resin
(WXDIX-4051), diglycidyl ether of bisphenol A, was
kindly provided by BLUESTAR Wuxi Petrochemical
Co., Ltd, Wuxi, China. This epoxy resin has an epoxide
equivalent weight of 880–950 g/mol and a soft point of
93–101
C. The direct-drawn glass fiber roving
(ECT4303) was purchased from Chongqing Polycamp
International, Chongqing, China. This product has a
linear density of 2400 tex and a filament diameter of
14 mm, and it has been surface treated with (3-glycidox-
ypropyl)trimethoxysilane coupling agent.
Processing of LFT composites
The raw PC and PBT resins were dried at 85
C for 12 h
to ensure that the moisture content is low enough to
prevent degradation. The pultrusion equipment used in
this work was custom designed and consisted of a fiber
creel, a preheating chamber, an impregnation appar-
atus coupled with a corotating twin-screw extruder
(d ¼ 30 mm, L/D ¼ 40, Nanjing Rubber & Plastics
Machinery Plant Co., Ltd, China), a wind cooler, a
pulling machine, and a custom-tailored pelletizer as
illustrated in Scheme 1. The preheating chamber is
used to preheat the glass fiber before it enters the
impregnation apparatus. The impregnation apparatus
consists of multiple heating zones, five rollers, a fiber
entrance nozzle, and a pultrusion die. When the glass
fibers are pulled from the fiber creel and induced into
the impregnation apparatus from the entrance nozzle
by the pulling machine, the fiber bundles will pass
through these rollers and achieve a good wet-out
effect. The pultrusion process is driven by a pulling
machine consisting of four pairs of polyurethane-
coated metal wheels. The pulling machine has been
designed for the requirements of force and speed for
the current experiment. On the other hand, the dried
PC and PBT pellets were premixed at different mass
ratios through a high-speed mixing machine, and mean-
while, the epoxy resin (1.5 wt% of total PC/PBT
amount) was added as a compatibilizing agent. Then
the PC/PBT blending melts were prepared through the
twin-screw extruder with a screw configuration adapted
to the blending of semicrystalline and amorphous poly-
mers. The temperatures along the barrel from feeding
zone to die were set at 235, 245, 260, 260, 260, 265, 265,
and 270
C, and the rotating speed of the screw was 60 r/
min. The continuous glass fiber was fed into the
impregnation apparatus including a number of thermo-
couples and roller, and meanwhile, the PC/PBT blend-
ing melt was delivered into the impregnation apparatus.
It should be mentioned that this impregnation appar-
atus was equipped with several replaceable exit nozzles
in different diameters. In this case, the amount of the
Scheme 1. Schematic preparation of the LFT composites through the melt pultrusion of glass fibers with PC/PBT allows.
1806 Journal of Reinforced Plastics and Composites 34(21)
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PC/PBT blending melt coated on glass fiber could be
well controlled when the impregnated fibers passed
through the given exit nozzles. Therefore, the fiber con-
tent of the composites could be accurately determined.
The continuous glass fiber impregnated with the PC/
PBT blending melt was pulled toward the custom-tai-
lored granulator after a wind cooling and then was cut
into the granules at a length of 10–12 mm. The tem-
perature in the impregnation unit was set to 275–
280
C, and the pulling speed of the fiber was adjusted
in the range of 6–12 m/min. The resulting granules were
dried under vacuum at 80
C overnight and then were
stored in the sealed aluminum foil bags.
Characterization
Measurement of mechanical properties. The LFT granule
samples were dried at 95
C in a vacuum oven for 6 h
and then were injection molded into testing bars using
an injection-molding machine (Haitian Machinery Co.
Ltd, Ningbo, China). The temperatures were set at 270,
275, 280, 280
C, injection pressure was 75 MPa, and the
rotating speed of the screw was 60 r/min. The tensile
and flexural properties of the injection-molded speci-
mens were measured using a computer-controlled
SANS CMT-4104 universal testing machine with a
load cell of 10 kN capacity in accordance with ISO-
527 and ISO-178 standards, respectively. The notched
Izod impact tests were performed on a SANS ZBC-
1400A impact tester equipped with a pendulum of
2.75 J according to ISO-180 standard. The impact test
specimens were notched on a milling machine to
achieve a ‘‘V’’ type notch in a depth of 2 0.2 mm.
All the measurements were conducted at room tem-
perature, and the reported values reflected a mean
value of five tests.
Fiber length analysis. An FASEP 3E-ECO system
(Emmeram Karg Industrietechnik, Germany) coupled
with scanner was used to determine the residual fiber
length and distribution of the injection-molded speci-
mens made from the LFT composites. The injection-
molded specimens were first calcined in the furnace at a
temperature about 800
C for 5 h to remove the PC/PBT
matrix. The fibers were released from their matrix and
then were transferred into ethanol on a glass slide.
Using the FASEP 3E-ECO system, the dispersed
fibers were counted and measured to achieve the resi-
dual length and distribution. The number average
length (
L
n
) could be calculated by the following
equation
L
n
¼
P
i
N
i
L
i
P
i
N
i
ð1Þ
Scanning electron microscopy (SEM). The fractography of
the LFT composites was performed on a Hitachi H-
4700 scanning electron microscope. The fracture sur-
faces obtained from the impact and tensile specimens
after impact and tensile measurements were made elec-
trically conductive by sputter coating with a thin layer
of gold–palladium alloy. The micrographs of SEM
were taken in high vacuum mode with 20 kV acceler-
ation voltage and in a medium spot size.
Differential scanning calorimetry (DSC). DSCs were per-
formed to determine the melting and crystallization
behaviors of the LFT composites on a TA instruments
Q20 differential scanning calorimeter equipped with a
Universal Analysis 2000 data station. All operations
were carried out under a nitrogen flow of 50 ml/min
and with a sample weight around 2–5 mg. During the
measurement, the samples should first be heated to
260
C and held at this temperature for 5 min to elim-
inate the effect of the thermal and processing history.
The crystallization and melting enthalpies (H
c
and
H
m
) were deduced by running a curve integral. The
degree of crystallinity (V
c
) was determined from DSC
scans using the following equation
X
c
¼
H
m
ð1 ’
1
’
2
ÞH
0
m
100% ð2Þ
where ’
1
and ’
2
are the weight percents of PC and
glass fibers, respectively, and H
0
m
is the melting
enthalpy of PBT in a 100% crystalline form, which
was set to 140 J/g.
39
Measurement of heat distortion temperature (HDT). The test
specimens for HDT were prepared through the injec-
tion-molding process. The HDTs of PC/PBT alloys and
their LFT composites were measured with a SANS
ZWK-1000 Heat distortion & Vicat softening tempera-
ture instrument under a load of 1.82 MPa according to
ISO-75A standard, and all the HDTs represented the
average values over five tests.
Dynamic mechanical analysis. The dynamic mechanical
behaviors of the LFT composites were measured on a
TA Q800 dynamic mechanical analyzer under a dual
cantilever mode. The specimens were cut out with a
dimension of 2 mm 10 mm 50 mm. The measure-
ment was run during the temperature range from 40
to 180
C with a heating rate of 5
C/min and an oscil-
lation amplitude of 10 mm at a frequency of 1 Hz.
Thermogravimetric analysis (TGA). TGA was performed on
a TA Instruments Q50 thermal gravimetric analyzer in
a nitrogen atmosphere. Samples were placed in a plat-
inum crucible and ramped from room temperature to
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750
C at a heating rate of 10
C/min, maintaining a flow
of nitrogen at 50 ml/min.
Results and discussion
Mechanical properties and fiber length distribution
The mechanical properties of the PC/PBT alloys and
their LFT composites with different loadings of glass
fibers were systematically evaluated in terms of the ten-
sile, flexural, and impact properties, and all of the test
results are summarized in Table 1. It is found that the
introduction of long glass fibers generates a significant
reinforcement effect on the PC/PBT alloys. As shown in
Table 1, the tensile strength of the PC/PBT composites
containing 20 wt% glass fibers was improved by 56% in
comparison to the PC/PBT alloy at a weight ratio of
70/30. Furthermore, the tensile strength increased by
160% with increasing the fiber loading up to 40 wt%.
Analog trend was also obtained in flexural strength and
modulus as shown in Table 1. Both the flexural strength
and modulus could be increased of 200% when 40 wt%
of glass fibers was incorporated. It is understandable
that the higher fiber loading could generate a denser
fiber framework to undergo structural stress in the
matrix, thus resulting in much higher strength and
modulus for the composites. It is noteworthy that
these mechanical properties are not only dependent
on the fiber loading but also relative to the weight
ratio of PC/PBT. In general, the higher the relative per-
centage of PC within the alloys, the better is the stiff-
ness. This may facilitate higher tensile and flexural
strength for the resulting LFT composites.
As reported with respect to the fiber-reinforced
thermoplastic composites, there are several factors
determining the reinforcement effect of fibers, which
include fiber content, fiber strength, fiber length, fiber
orientation, interface adhesion, and polymeric matrix
nature.
16,18,40
It is clear that the pultruded long glass
fibers through the melt of PC/PBT alloys can deliver
significantly longer fiber to the injection-molded com-
posites in comparison to the extruded short fiber ones.
The aspect ratio (ratio of fiber length (l) to diameter ( d ))
of fibers in the resulting LFT composites is an order of
magnitude greater than that of the short fibers, often
exceeding l/d of 2000 and, thus, takes full advantage of
the strength of the reinforcing fiber.
16
There is no doubt
that the residual fiber length is considered as a critical
factor determining the reinforcement effect of the
resulting composites as a result of the application of
LFT processing technique. Therefore, it is of great sig-
nificance to understand the residual fiber length and its
distribution within the injection-molded specimens of
the PC/PBT composites. Figure 1 shows some
Table 1. The mechanical properties of the LFT composites based on PC/PBT alloys.
Sample composition (wt%)
L
n
(mm)
Tensile
strength (MPa)
Elongation at
break (%)
Flexural
strength (MPa)
Flexural
modulus (MPa)
Notched
Izod impact
strength (kJ/m
2
)PC PBT Glass fibers
100 0 0 – 61.57 135.7 79.7 211.5 9.54
0 100 0 – 51.48 6.9 76.0 2050.5 2.84
40 60 0 – 50.58 131.3 78.2 2494.7 9.20
30 70 0 – 50.46 137.2 77.8 2158.0 7.81
20 80 0 – 50.52 140.2 75.9 2088.4 7.25
36 54 10 2.77 70.16 9.5 85.8 4125.5 10.32
27 63 10 2.81 75.44 10.9 94.0 4585.2 8.21
18 72 10 2.88 72.05 8.4 88.4 4103.2 8.07
32 48 20 2.72 96.63 11.2 150.5 5900.7 16.52
24 56 20 2.83 85.31 10.0 141.2 5692.1 15.85
16 64 20 2.85 78.37 10.8 112.3 4770.4 14.23
28 42 30 2.75 116.04 11.7 159.5 6786.3 17.92
21 49 30 2.75 111.1 11.5 184.9 7788.9 16.36
14 56 30 2.79 119.05 12.5 167.0 8024.7 16.05
24 36 40 2.55 137.44 12.3 210.9 11146 18.71
18 42 40 2.56 130.51 12.5 193.2 11043 18.46
12 48 40 2.58 128.72 12.4 205.8 9463.7 17.52
1808 Journal of Reinforced Plastics and Composites 34(21)
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representative distribution charts for the fiber length
within the injection-molded specimens, and meanwhile,
the corresponding optical micrographs are presented as
insets. Furthermore, the number average fiber lengths
were determined through the fiber length analysis soft-
ware equipped with FASEP system and are also sum-
marized in Table 1. It should be mentioned that, by
using the FASEP method of nondestructive ashing
glass-fiber-reinforced thermoplastic composites, the
fibers are extracted from work-piece without breaking.
It is visibly found from Figure 1 that most of the glass
fibers present a residual length ranging from 0.5 to
3.5 mm. Many studies indicated that, for the short-
fiber-reinforced composites, their residual fiber length
could usually be maintained in the range of 0.1–
0.5 mm.
40–43
It was reported that the average residual
length was lower than 0.3 mm when the PBT resin was
reinforced with short glass fibers.
44
Moreover, the data
listed in Table 1 demonstrate that the LFT composites
have a prominent superiority in the measured average
fiber length compared to the short-fiber-reinforced ones
as reported by Hashemi.
44
These results imply that the
residual fiber length was reduced significantly from 10
to 12 mm of the feedstock during the injection-molding
process due to a strong shear effect of injection screw at
the plasticization stage of the composites. However, the
fiber loading seems to influence the residual length, and
the higher fiber loading leads to a shorter residual
length. This is due to the enhanced shearing action
among the glass fibers. Nevertheless, the weight ratio
of PC/PBT seems to have no influence on the residual
fiber length within the thermoplastic matrix. In any
case, the final length of glass fibers within the specimens
is still much longer than that of the short-glass-fiber-
reinforced ones, in which the residual fiber length is
only distributed in the range of 0.1–0.5 mm.
40,41
Figure 1. Fiber length distribution charts of the LFT composites at PC/PBT/glass fibers weight ratios of (a) 24/56/20, (b) 18/42/40,
(c) 12/48/40, and (d) 24/36/40; inset: the corresponding optical micrographs of residual fibers.
Tan et al. 1809
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By determination of the residual fiber length and
fiber arrangement, a prediction for the mechanical
properties of fiber-reinforced composites is possible.
Based on the fiber pullout mechanisms for the fiber-
reinforced thermoplastic composites, the Kelly–Tyson
model established a relationship between the ultimate
tensile stress and the fiber length by the following
equation
45
uc
¼
X
l
i
5 l
c
l
i
V
i
d
f
þ
X
l
j
4 l
c
f
V
j
1
l
c
2l
j
2
4
3
5
þ
um
ð1 V
f
Þ
ð3Þ
where
uc
and
f
are the ultimate tensile strength of
composite and fiber, respectively;
um
the matrix
stress at the failure of the composite; V
f
the volume
fraction of fibers; the orientation factor; l
c
the critical
fiber length; d
f
the diameter of fiber; and is the inter-
facial shear stress between the fiber and matrix within
composites. In this equation, the contributions from the
glass fibers with the subcritical and supercritical fiber
lengths, and the resin matrix are clearly distinguished.
For the LFT composites based on the PC/PBT alloys,
the V
f
could be calculated by means of equation (4)
17
V
f
¼ 1 þ
f
m
1 W
f
W
f
ð4Þ
where W
f
is the weight fraction of glass fibers,
m
the
density of the PC/PBT matrix, and
f
is the density of
glass fibers (2.54 g/cm
3
). The fiber length (l
i
) could be
directly obtained from the optical FASEP images of the
fibers extracted from the specimens. The orientation
factor was set to 0.5 for the injection-molded specimens
with the free flow.
17
Meanwhile, according to the
Kelly–Tyson model, the critical fiber length could be
calculated by the following equation
17,46
l
c
¼
f
d
f
2
¼
E
f
"
f
d
f
2
ð5Þ
where E
f
is the elastic modulus of fibers, and "
f
is the
strain at break of fibers. If the is set to the shear stress
of thermoplastic matrix as the upper limit and E
f
is
taken to be a constant of 74 GPa for glass fibers, the
l
c
is calculated to be 0.618–0.674 mm on the basis of the
mechanical results of the LFT composites. In this case,
the predicted tensile stress for the LFT PC/PBT system
could be calculated by equations (3) to (5). Figure 2
shows the theoretical ultimate tensile stress of the
LFT composites as a function of the fiber loading in
the case of PC/PBT (30/70) weight ratio, and the cor-
responding experimental data were presented simultan-
eously. It is interesting to note that, in principle, the
theoretical results are similar to the measured ultimate
tensile stress and show the same trend as a function of
fiber concentration. It is evident that the average fiber
length is significantly greater than the critical fiber
length, and thus, these longer fibers can contribute
much better reinforcement effect for the resulting com-
posites. Such a contribution from the fibers with a
supercritical fiber length is superior to those with the
subcritical one and thus results in a much higher rein-
forcing efficiency. Nevertheless, the theoretical data
seem to be higher than the experimental one as shown
in Figure 2. The interfacial shear stress between the
matrix and glass fibers is assuredly lower than the
shear stress among the matrix itself, and therefore,
the thermoplastic matrix requires a much longer fiber
length than the critical one to achieve an expected
reinforcement effect. This may lead to a lower measured
tensile stress for the LFT composites.
The impact toughness of the LFT composites was
also measured, and the results are summarized in
Table 1. These impact data were found to present a
similar trend to the results previously observed for ten-
sile and flexural behaviors. The impact strength of the
composites shows a considerable improvement with
increasing the fiber loading, and meanwhile, it achieves
an increment by a factor of two over the whole ratios of
PC/PBT for the thermoplastic matrix. Such a notable
result is attributed to the reinforcement effectiveness
derived from long glass fibers. However, for most of
the composites with short reinforcing fibers, the
impact strength only shows an improvement by 10–
20% in comparison to their pure resins due to the dis-
tribution of the residual in the range of 0.1–0.5 mm as
reported by the corresponding literature.
40–42,47,48
Figure 2. Theoretical and measured tensile stress as a function
of fiber loading for the LFT composites at a PC/PBT weight ratio
of 30/70.
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The high levels of impact strength exhibited by the LFT
composites are often attributed to energy dissipating
mechanisms of fiber debonding, pullout, and fracture,
which have been discussed by many researchers.
18–21
The impact energy dissipation is also generally influ-
enced by several factors including matrix fraction,
fiber debonding, friction between the interfaces of
matrix and fibers, fiber pullout, and fiber breakage.
49
Like the Kelly–Tyson model for the tensile stress of the
composites, the Cottrell model is established with two
parts depending on whether the fiber length is subcrit-
ical or supercritical. This model can predict the impact
energy (U
c
) of a fiber-reinforced composite through the
following equations
50,51
U
c
¼ U
m
ð1 V
f
Þþ
V
f
lU
d
d
f
þ
V
f
l
2
f
6d
f
, when l 5 l
c
ð6Þ
U
c
¼ U
m
ð1 V
f
Þþ
V
f
ðl l
c
ÞU
f
l
þ
V
f
l
2
c
U
d
d
f
l
þ
V
f
l
3
c
f
6d
f
l
, when l 4 l
c
ð7Þ
where U
m
, U
f
, and U
d
are the fracture energy for
matrix, fibers, and interfaces, respectively, and
f
is
the interfacial friction during fiber pullout. For the
composite systems with a fiber length lower or close
to the critical one like short-fiber-reinforced thermo-
plastic composites, the impact energy dissipation
should be estimated by equation (6), where the three
terms cover matrix fracture, fiber–matrix debonding,
and fiber pullout. The fiber pullout is the most import-
ant energy dissipation mechanism. Considering of the
LFT composites with a supercritical fiber length, the
impact energy dissipation should mainly be predicted
by equation (7), where the four terms account for
matrix fracture, fiber–matrix debonding, fiber fracture,
and fiber pullout limited to the critical length. The con-
tribution of energy dissipation from matrix is typically
limited in contrast to the reinforcement, because the
energy dissipation by the matrix is small due to the
fact that the presence of fibers prevents large deform-
ation of the matrix. The fiber pullout seems to be one of
the predominant energy absorption mechanisms. On
the other hand, most of the fibers dispersed in the
matrix have a supercritical length, so the fiber breakage
is inevitable to occur as well as the fiber pullout, thus
dissipating a great amount of impact energy. However,
there are still several uncertain critical parameters in
equation (7), and therefore, the prediction for the
impart energy is unavailable in the current work. In
any case, the long glass fiber not only supplies promin-
ent reinforcement toward the PC/PBT alloys but also
evidently gives its greatest benefits to the impact tough-
ness. In the following section, the SEM investigation on
the fractography will provide more evidence for the
impact and tensile behaviors of the LFT composites.
Morphology and microstructures
The SEM observation was performed to investigate the
tensile fracture surfaces derived from the tested tensile
specimens of the LFT PC/PBT composites, and the
obtained micrographs are illustrated in Figure 3.
These SEM micrographs clearly indicated some
detailed information about fiber orientation and inter-
facial adhesion between the glass fibers and resin
matrix. As shown in Figure 3, a large number of
pulled-out fibers as well as some hollows on the cross-
sectional areas are distinctly observed. Most of the
fibers on the fracture surfaces are found to have a
length over the aforementioned critical length for the
LFT composites. Meanwhile, the fiber breakage is also
observed on the fracture surfaces. These results confirm
that both the fiber pullout and fiber breakage are pre-
dominant for the reinforcement of the LFT composites
based on the PC/PBT alloys. It is noteworthy that the
fiber surfaces are very coarse with obvious resin matrix
sticking on them when the SEM micrographs are
viewed with a higher magnification (see Figure 3(d),
(f), and (h)), which indicates a good interfacial adhesion
between the fibers and matrix. Such a high adhesion
can significantly enhance the stress transfer when the
fiber debonding occurs through frictional forces along
the interfaces. These results further confirm that the
long glass fibers have a much stronger reinforcing cap-
ability toward thermoplastics than the short ones, thus
causing more significant reinforcement efficiency for the
PC/PBT matrix. On the other hand, the impact fracture
surfaces of the LFP composite specimens were also
investigated with SEM. Figure 4 shows some represen-
tative SEM micrographs of PC/PBT alloys and their
LFT composites. It is observed from Figure 4(a) that
the PC/PBT alloy exhibits a considerably coarse frac-
ture surface due to the significant matrix deformation
when the impact fracture occurs. There is no visible
interface failure observed in the fracture surface, indi-
cating good compatibility between the two phases of
the alloy. Similar to the tensile fracture surfaces, the
SEM micrographs of the impact fracture surfaces of
the LFT composites are also found to have a feature
of major fiber pullout, and meanwhile, a few of fiber
section planes due to breakage are observed. It is not-
able from Figure 4(b) to (f), compared to the pristine
PC/PBT alloy (see Figure 4(a)), the LFT composites
matrix did not show more significant matrix deform-
ation on the impact fracture surfaces. These results fur-
ther confirm the predominant contribution of the fiber
Tan et al. 1811
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pullout and fiber breakage due to the dissipation of
impact energy. Moreover, the remaining fibers on the
fracture surfaces are observed to be rough with some
sticking resins due to the fiber–matrix debonding, indi-
cating a high level of interfacial adhesion between the
fibers and matrix. These phenomena suggest that a
mechanical interlocking has been established between
the fibers and matrix, and thus, the better stress transfer
can be gained. The glass fibers may fracture if the fiber
stress level exceeds the local fiber strength. The fibers
that have fractured away from the crack interface are
pulled out of the matrix, which may also be involved in
energy dissipation. Therefore, the PC/PBT alloys rein-
forced with long glass fiber exhibit much better impact
resistance than those with the conventional
short one.
21,41
Figure 3. SEM micrographs of tensile fracture surfaces for the LFT composites at PC/PBT/glass fibers weight ratios of (a) 24/56/20,
(b) 21/49/30, (c, d) 18/42/40, (e, f) 12/48/40, and (g, h) 24/36/40.
1812 Journal of Reinforced Plastics and Composites 34(21)
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Thermal behaviors
The LFT composites developed in this study is a type of
PBT-rich thermoplastic composites and mainly consist
of three components, i.e. semicrystalline PBT, amorph-
ous PC, and inorganic glass fibers. Their properties and
processability for the final products are considerably
associated with the melting and crystallization charac-
teristics of the PBT phase within the composites.
Meanwhile, the crystallization behaviors of the PBT
phase are significantly influenced by both the PC
phase and glass fibers, which may result in a complica-
tion of crystallinity of PBT phase in most cases.
Therefore, it is of great importance to investigate the
melting and crystallization behaviors of the PBT phase
within the LFT composites. The dynamic DSC scans
were first performed to investigate the crystallization
behaviors of the composites. Figure 5 shows the DSC
cooling thermograms of some representative PC/PBT
alloys and their LFT composites. The corresponding
melting processes were also recorded by the dynamic
DSC scan and are illustrated in Figure 6. It is observed
that the pure PBT exhibits a single crystallization peak
on its DSC curve, whereas the PC/PBT alloys present a
multiple crystallization behavior. From the DSC curves
of the PC/PBT alloys as shown in Figure 5, the same
crystallization exothermic peak at a low temperature
around 125
C could be observed together with the
main crystallization peak as well as a shoulder corres-
ponding to formation of the b-crystal. This phenom-
enon was also reported by Delimony et al. and was
ascribed to fact that the crystallization of the residual
amorphous PBT within the continuous crystalline PBT-
rich phase occurred at high supercoolings.
52
Such a
multiple crystallization behavior resulted in the mul-
tiple melting features for the PC/PBT alloys as depicted
Figure 4. SEM micrographs of impact fracture surfaces of the LFT composites at PC/PBT/glass fibers weight ratios of (a) 30/70/0, (b)
24/56/20, (c) 21/49/30, (d) 18/42/40, (e) 12/48/40, and (f) 24/36/40.
Tan et al. 1813
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by Figure 6. It should be highlighted that both the PC
phase and glass fibers affected the main crystallization
temperature (T
c
) of PBT phase within the LFT com-
posites as observed from the thermal analysis data in
Table 2. The formation of PC/PBT alloys evidently
reduced the T
c
of PBT phase due to the interference
on the PBT crystallization from the PC macromol-
ecules. However, the introduction of glass fibers
seems to increase the T
c
slightly, especially in high
fiber loadings. This is attributed to the fact that the
heterogeneous nucleating effect of glass fibers facilitates
the nucleation and growth of spherulites and thus
enhances crystallinity of the PBT phase. Furthermore,
from the crystalline data listed in Table 2, it is noted
that the crystallinity degree of PBT phase was improved
considerably in the presence of glass fibers, although
the growth of its spherulites was suppressed due to
the incorporation of PC. Moreover, the enhancement
of crystallinity also diminished the unexpected low-tem-
perature crystallization of PBT phase within the com-
posites and improved the processability of the LFT
composites accordingly.
As one of the most important thermal properties of
engineering materials, the elevated temperature per-
formance reflects the thermal resistance and thermal
dimensional stability of a material and, therefore,
establishes the service environment for the structural
engineering materials. The HDT is defined as the
upper limiting temperature of the dimensional stability
of materials in service without significant physical
deformations under a normal load and thermal effect.
In general, the HDT can be used to estimate the ele-
vated temperature performance.
53
The HDTs of the
PC/PBT alloys and their LFT composites are listed in
Table 2. The pure PBT is found to only have an HDT
of 55.6
C. However, the PC/PBT alloys also show the
HDTs as low as pure PBT, indicating that the dispersed
PC phase could not make a positive effect on the HDT
of the PBT-rich phase. Such a low HDT cannot meet
the requirement for the PC/PBT alloy as an engineering
material in most industrial and civil applications. It is
surprisingly noted in Table 2 that the HDT jumped to
161
C only by introducing 10 wt% glass fibers and then
continued to increase up to 204
C at a fiber loading of
40 wt%. It is well known that the HDT is commonly
influenced by several factors such as the composition,
the additives like nucleating agents or fillers, and pro-
cessing conditions. Moreover, it is associated with the
mechanical behavior of thermoplastic composites.
Nielsen and Landel
54
proposed that the HDT was rela-
tive to the flexural behaviors, especially the flexural
modulus as a result of thermoplastics compounding
with inorganic fillers. It is understandable that, in the
current work, the presence of long glass fibers restricts
the deformation of thermoplastic matrix due to the for-
mation of the LFT composites and thus increases the
modulus of the resulting composites. Such a trend is
increasingly significant with increasing the fiber load-
ing. In this case, the LFT composites could retain high
stiffness and high modulus with increasing temperature
and achieve high HDTs accordingly. The results
obtained in this study are in good agreement with the
Nielsen’s prediction. In addition, the improvement in
HDT may also be due to the fact that glass fibers
Figure 6. DSC heating thermograms of the LFT composites at
different weight ratios of PC/PBT/glass fibers (GF).
Figure 5. DSC cooling thermograms of the LFT composites at
different weight ratios of PC/PBT/glass fibers (GF).
1814 Journal of Reinforced Plastics and Composites 34(21)
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promote the crystallization of PBT phase within the
composites and thus inhibit the deformation of the
composites.
19
Dynamic mechanical analysis
The dynamic mechanical properties of the PC/PBT
alloys and their LFT composites were investigated by
DMA. The resulting storage modulus and loss factor
(tan d) as a function of temperature are illustrated in
Figures 7 and 8, respectively. The dynamic mechanical
behaviors of pure PC and PBT resins are also presented
as references in the two figures. It is observed from
Figure 7 that the PC resin shows a slightly higher stor-
age modulus than PBT in the temperature region below
glass transition temperature (T
g
). However, the storage
moduli of these two resins present a sharp decline in the
glass transition zone, which was confirmed by the loss-
factor peaks corresponding to the temperature depend-
ence of tan d as shown in Figure 8. The PC/PBT alloys
seem to keep the similar level of storage modulus with
PBT before the glass transition occurs. Nevertheless,
the PC/PBT alloys show two relaxation peaks assigned
to the glass transitions of the PBT-rich phase and the
dispersed PC phase as observed in Figure 8. This phe-
nomenon implies that the PBT phase is thermodynam-
ically immiscible with the PC one within the alloys.
However, the glass transition peak of the PC phase
shows a significant shift to a low temperature,
indicating a significant change in the macromolecular
structure and chain length of the PC phase. This may be
attributed to the transesterification occurring between
the PC and PBT phases as well as the thermal degrad-
ation of PC, because the two polymer melts had to stay
in the hot chamber of the impregnation apparatus for a
long time during the pultrusion process. The transester-
ification of PC with PBT made a connection between
two polymeric chains, thus resulting in partial miscibil-
ity between the PC and PBT.
29,30
Moreover, the add-
ition of epoxy resin as a compatibilizer also enhanced
the compatibility of two phases in the alloys. These
effects collaboratively lead to a remarkable shift of
relaxation peak of the PC phase.
55,56
It is important
to note that the introduction of glass fibers significantly
improved the storage modulus of the resulting compos-
ites in all temperature ranges, especially at the tempera-
tures lower than the T
g
of the PBT phase, and
meanwhile, the storage modulus could be further
improved with increasing the fiber loading. These
results are in good agreement with the static tensile
and flexure moduli of these composites. It is believable
that the excellent reinforcing and stiffening effect of
long glass fibers toward the thermoplastic matrix as
well as the physical interaction between the fibers and
matrix may jointly contribute an improvement in resili-
ence to the composites. Moreover, the LFT composites
with lower weigh ratios of PC/PBT are found to present
higher storage modulus at the same fiber loading. This
Table 2. The thermal properties of the LFT composites based on PC/PBT alloys.
Sample composition (wt%) Heat deflection
temperature
(
C, 1.82 MPa)
DSC data TGA data
PC PBT Glass fibers T
m
(
C) X
c
(%) T
onset
(
C) T
max
(
C) Char yield (wt%)
100 0 0 135.5 – – 367.3 412.3 21.38
0 100 0 59.9 190.83 44.5 379.8 396.5 5.97
40 60 0 60.2 177.82 41.2 381.4 409.2 13.54
30 70 0 59.3 178.23 42.6 380.1 406.8 11.12
20 80 0 59.0 179.45 44.2 376.2 396.5 10.78
36 54 10 161.4 177.86 41.6 382.5 410.4 15.49
27 63 10 161.7 177.95 42.1 382.7 409.6 13.62
18 72 10 162.3 176.54 42.8 380.5 405.2 13.74
32 48 20 174.6 176.30 42.1 384.7 411.8 27.53
24 56 20 172.8 176.42 42.6 382.9 409.3 25.28
16 64 20 175.1 176.86 43.4 383.2 407.2 23.87
28 42 30 196.8 177.16 42.5 384.5 415.4 35.24
21 49 30 189.4 177.85 43.8 383.6 412.7 32.46
14 56 30 194.9 180.55 44.5 382.3 410.6 33.19
24 36 40 204.9 178.25 44.2 385.2 415.8 45.71
18 42 40 202.3 178.93 45.3 383.4 413.5 44.62
12 48 40 203.9 180.79 46.5 383.9 412.1 42.58
Tan et al. 1815
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indicates that the new interfaces within the composites
are disadvantageous to the mechanical resilience.
In addition, it is noteworthy in Figure 8 that, for all
of the alloys and their LFT composites, the loss-factor
peaks show slight shifts to higher or lower temperatures
around the glass transition peak of the PBT-rich phase
as a function of composition. This indicates the com-
plicated effects of miscibility, intermolecular inter-
action, and interfacial feature on the macromolecular
motility due to the formation of multiphase within the
composites. Furthermore, it is also noteworthy that, for
all of the composites, the magnitude of glass transition
peak corresponding to the PBT-rich phase is higher
than that of pure PBT but lower than those of alloys.
It is well known that the damping in the glass transition
zone measures the imperfection in the elasticity and
that much of the energy expended for the deformation
of a material during DMA testing is dissipated directly
into heat.
57
This means that the macromolecular mobil-
ity within the composites decreases with introduction of
glass fibers, which results in a reduction of mechanical
loss to overcome inter-friction between molecular
chains. Generally, the damping of the polymer is
much greater than that of the fibers. The incorporation
of inorganic fibers into polymer matrix can increase its
elasticity and reduce its viscosity, and thus less energy is
consumed to overcome the friction forces between the
molecular chains.
57
As a result, the loss-factor magni-
tudes of the composites decrease in the presence of glass
fibers.
TGA
The thermal degradation behaviors and thermal stabi-
lities of the PC/PBT alloys and their LFT composites
were evaluated by TGA under a nitrogen atmosphere.
Figure 9 shows the resulting thermograms of some
representative samples, and all of the TGA data are
summarized in Table 2. It is interesting to note from
these weight-loss profiles that both pure PBT and PC
present a typical one-step thermal degradation behav-
ior in nitrogen. However, the PC/PBT alloys and their
Figure 7. Temperature dependence of storage modulus from the DMA measurement for the LFT composites at different weight
ratios of PC/PBT/glass fibers (GF).
1816 Journal of Reinforced Plastics and Composites 34(21)
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composites show a weak second-stage decomposition at
higher temperatures besides a main degradation, which
is due to the fact that the PC phase has a higher decom-
position temperature than the PBT phase. In this case,
the first weight loss is ascribed to the decomposition of
the PBT-rich phase, whose rapid decomposition along
with a maximum weight loss occurs due to the random
scission of PBT main chains during the pyrolysis pro-
cess. The second weight loss can be assigned to the
decomposition of the dispersed PC phase occurring in
a wide temperature region. The two characteristic tem-
peratures corresponding to the 5 wt% weight loss and
maximum weight loss rate are defined as the onset
decomposition temperature (T
onset
) and maximum
decomposition temperature (T
max
), respectively, and
both of them are furthermore taken as indicators of
thermal stability. From Table 2, it is clearly observed
that the T
onset
’s and T
max
’s of the PC/PBT alloys are
higher than those of PBT but lower than those of PC.
Figure 8. Temperature dependence of loss factor (tan ) from the DMA measurement for the LFT composites at different weight
ratios of PC/PBT/glass fibers (GF).
Figure 9. TGA thermograms of the LFT composites at differ-
ent weight ratios of PC/PBT/glass fibers (GF).
Tan et al. 1817
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Nevertheless, the PC/PBT alloys show a slight rise in
T
max
with increasing the proportion of the more ther-
mally stable PC within the alloys. It is important to note
an improvement in T
max
with introducing glass fibers
into the PC/PBT alloys, indicating the long glass fibers
seem to enhance the thermal stability of the composites
slightly, though there is no influence on thermal decom-
position behaviors of the thermoplastic matrix. Some
studies reported that the presence of inorganic fillers
could reduce the heat release rate of the polymeric
matrix, thus retarding the decomposition of the polymer
domain.
58
Moreover, the glass fibers could also act as
physical barriers to hinder the transport of volatile
decomposed products out of the LFT composites
during thermal decomposition. Therefore, the thermal
stabilities of the composites show an increasing trend
with increasing the fiber loading. In addition, the char
yields seem to increase at a level higher than the weight
percentages of glass fibers in the composites, indicating
that glass fibers effectively enhance the char formation
of the thermoplastic matrix. In general, the glass fibers
can generate a heat barrier to thermal decomposition
products and enhance the carbonization of the thermo-
plastic matrix during the thermal degradation process,
and consequently, the higher char yields are achieved at
the end of the thermal decomposition of the composites.
Conclusion
The LFT composites based on the PC/PBT matrix with
different loadings of glass fibers were successfully pre-
pared through the melt blending followed with melt
pultrusion using the custom-designed pultrusion equip-
ment to manufacture the thermoplastic long granules.
The resulting composite specimens revealed excellent
mechanical properties due to the superiority of the
LFT technique. The levels of improvement in tensile
strength, flexural strength and modulus, and impact
toughness are dependent on the fiber loading within
the composites. Such a prominent reinforcement effect
is attributed to the feature that the residual fiber length
in the injection-molded LFT composite specimens is
greatly superior to that in the short-fiber-reinforced
ones. This takes full advantage of the strength of the
reinforcing fibers themselves. The enhancement of
impact toughness is due to the energy dissipation by
both the fiber pullout and fiber fracture as a result of
the long-fiber-reinforcement effectiveness. The morpho-
logic observation confirmed that the fiber pullout and
fiber breakage concurred on the impact and tensile frac-
ture surfaces, and the former was more significant than
the latter. The SEM investigation also indicated a good
interfacial adhesion between the thermoplastic matrix
and fibers, which provided subsidiary mechanical
enhancement for the composites. The LFT composites
also obtained a remarkable increase in storage modulus
but presented a considerable decrease in loss-factor
magnitude compared to the pristine alloys. In addition,
the thermal stabilities of the composites were slightly
improved in the presence of long glass fibers due to the
heat barrier effect of inorganic fillers.
Declaration of conflicting interest
The author(s) declared no potential conflicts of interest with
respect to the research, authorship, and/or publication of this
article.
Funding
The author(s) declared the following potential conflicts of
interest with respect to the research, authorship, and/or pub-
lication of this article: The financial support from the
National Natural Science Foundation of China (Project
Grant no.: 51173010) is also gratefully acknowledged.
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