Fracture toughness and impact behaviour of glass-fibre-reinforced polyamide 6,6 injection mouldings

E L S E V I E R 0 2 6 6 - 3 5 3 8 ( 9 5 ) 0 0 0 7 5 - 5

Composites Science and Technology 55 (1995) 109-118 (~) 1995 Elsevier Science Limited

Printed in Northern Ireland. All rights reserved 0266-3538/95/$09.50

FRACTURE TOUGHNESS AND IMPACT B E H A V I O U R OF GLASS-FIBRE-REINFORCED POLYAMIDE 6,6

INJECTION MOULDINGS

M. Akay, D. F. O'Regan* & R. S. Bailey; School of Electrical and Mechanical Engineering, University of Ulster at Jordanstown, Newtownabbey BT37 OQB, UK

(Received 29 July 1994; revised version received 1 May 1995; accepted 5 May 1995)

Abstract The fracture toughness o f injection-moulded long-glass- fibre-reinforced polyamide 6,6 (average fibre aspect ratio in mouldings, l/d -~ 55) with fibre weight contents o f 50 and 60% was studied and compared with that o f short-glass-fibre-reinforced polyamide 6,6 (l/d ~ 20). Toughness was measured by using Mode I compact- tension specimens cut from the mouldings and also by means o f instrumented drop-weight impact tests. The addition of fibres to a polyamide 6,6 matrix led to an improvement o f fracture toughness. The improvement was marginally greater in the long-fibre grades. Fracture toughness values were dictated by the average in-plane fibre orientation angle ahead of the crack front and fibre length, particularly for the long-fibre grades.

Instrumented drop-weight impact tests indicated brittle fracture for unreinforced polyamide 6,6, causing possible catastrophic shattering into several pieces. Short-fibre and long-fibre materials both remained intact after impact, with smaller dimensions of the damaged area for the long-fibre materials. Crack initiation force and total energy values were increased by 25-37% and 300-600%, respectively, with the inclusion o f glass fibres. Negligible differences in these values were observed between the long- and short-fibre grades and with the long-fibre content increasing from 50 to 60% w/w.

Keywords: polyamide 6,6, glass fibre, fracture toughness, impact properties, fracture surface, fibre length, fibre orientation

1 INTRODUCTION

Discontinuous-fibre-reinforced thermoplastic compounds are commercially important in engineering applications. These materials may be processed by conventional methods, such as injection moulding, and

* Present address: A. Schulman GmbH, Htittenstrage 211, Kerpen-Sindorf, Germany. :~ Present address: ICI Wilton Centre, Wilton, Cleveland TS6 8JE, UK.

offer improvements in physical properties over those of unfilled materials. They replace metals in various applications because of their ease of fabrication, light weight and economy. However, additional considera- tions include processing-induced anisotropy in terms of through-thickness fibre orientation and length within moulded parts. Material defects such as voids and cracks may be also present or initiated in one of three regions--the matrix, the fibres, or the fibre/matrix interface and influence the load bearing capacity of the finished component.

The mechanisms of failure of short-fibre composites are now well documented) -9 Early work by Blumentritt et al. 1 using different fibre types (e.g. E-glass and graphite) and polymer matrices (e.g. polyethylene and polyamide 12), found fracture to occur by fibre pull-out rather than breakage, independent of whether brittle or ductile matrices were used. However, later work by Mandell et al., 2 investigating modes of crack propagation in injection- moulded short-glass and carbon-fibre-reinforced thermoplastics, established that the main crack will grow in a fibre avoidance mode (i.e. a path of least resistance) and the local mode of crack tip advancement varies with matrix ductility and the interfacial bond strength. Similar findings were reported by others. 3'4

Sato et al. s describe the mechanism of fracture of similar materials as: (i) initiation of the interracial cracks at fibre ends; (ii) propagation of such cracks along fibre sides; and (iii) crack propagation into the matrix leading to the failure of the material. In a later work Sato et al. 6 went on to describe the microfailure proceeding in five steps, consisting of the first two steps listed previously, followed by: (iii) the formation of plastic deformation bands in the matrix; (iv) crack opening within the band; and (v) a catastrophic crack propagation through the matrix with fibres being pulled out from the matrix

Sato et al. 6 also proposed some methods of improving composite toughness: (i) the use of fibres with small diameters ( ~ 10/xm), which reduce stress

109

110 M. Akay, D. F. O'Regan, R. S. Bailey

concentration at fibre tips due to the closer packing of the fibres; (ii) the use of longer fibres than conventional ones, these have the effect of reducing the number of fibre tips which act as stress concentrators; and (iii) the introduction of interfacial layers between the fibre and matrix. If a soft layer such as rubber could be used, the initiation of the interfacial microfailure would be expected to be reduced due to large deformation of the layer. By contrast, the use of a rigid layer with a modulus intermediate between that of the fibre and matrix, is considered to lower the stress concentration along the fibre sides.

Carling and Williams 7 investigated the effects of fibre length, and they found that there was little difference in fracture toughness when as-moulded specimens (dry) were tested. However, after being hot and wet conditioned, yielding 2-3% moisture absorption in the specimens, the longer fibre material showed significant improvements over the shorter fibre materials.

Leach and Moore 9 investigated the failure and fracture behaviour of a range of reinforced Nylon compounds. Instrumented falling weight impact tests have been found to provide a good resolution of initiation and propagation mechanisms. Total impact energy to fracture showed a decrease with fibre content up to approximately 15% w/w, thereafter it increased and reached a plateau at a fibre loading of approximately 33% w/w. Choice of matrix was also found to influence both the energy to initiate and propagate cracks. Unreinforced Nylon 6 appeared less tough than Nylon 6,6, although as glass content increased the variations almost disappeared. As would be expected greater energy absorption was exhibited in using toughened unfilled Nylon 6,6.

The aim of the present work was to examine the fracture toughness and impact behaviour of a range of glass-fibre-reinforced polyamides of high fibre content which were promoted as potential metal replacement materials.

2 EXPERIMENTAL

Power-take-off (PTO) covers for a truck and flat plaque mouldings (Fig. 1) were injection-moulded from polyamide (PA) based materials, as detailed in Table 1, using a microprocessor-controlled Battenfeld BA950/315CD and an Engel CC90. Machine settings were 91 MPa injection pressure, 80 rpm screw speed, 1 MPa back pressure, 47 MPa hold pressure, 11 s hold pressure time, a barrel temperature profile of 270-300°C and a mould temperature of 80-100°C. The raw materials were dried in a vacuum oven at 80°C for 24 h and cooled to room temperature under vacuum prior to moulding.

2.1 Fibre orientation The through-thickness fibre orientation was determined using image analysis techniques (Seescan Solitaire Image Analyser). Image analysis relies on reflective microscopy of polished surfaces, where the intersections of the fibres with the surfaces form elliptical images. Measurement of the shape (the major and minor axes) and the orientation of the ellipse for each image was used to compute fibre orientation angles (out-of-plane, fl, and in-plane, a). Samples were machined in the x - z plane (i.e. through-the-thickness parallel to the fill direction) from the mouldings in locations which approximated to sites ahead of the crack front. Each sample was mounted in cold cure epoxy, ground and polished on a Struers DAP7 Pedemin-2 automatic polishing machine to a final stage using 1/xm diamond spray.

2.2 Fibre length distributions Fibre length distributions (FLD) were determined using image analysis techniques. This required the isolation of the fibres from the composite material by burning the matrix at 600°C in a muffle furnace for 1 h, to leave an ash of fibrous material. Long-fibre mouldings produced a skeletal structure of fibres after burn off, with distinct surface and core layers of fibre orientations. Accordingly analysis of surface/core layers based on FLD may be carried out. Conversely, the short-fibre ash does not remain intact, hence separate surface/core FLD analysis was not at- tempted. In this method 0.03g of fibres were extracted from the sample ash and dispersed in water in a square glass dish. The dish was then placed on the observation stage of an Olympus SZH40 microscope and viewed in dark field transmission light. A video camera attached to the microscope transmitted live fibre images to the Seescan Solitaire Image Analyser. The dedicated software automatically digitised the fibre image and enhanced the contrast between foreground (fibres) and background. Measurements commenced on the calibrated system, when the touching and crossed fibres are automatically separ- ated. More than 1000 fibres were measured from each of the surface/core zones.

Fibre volume content at various positions of the mouldings, except through the thickness, was also determined by the process of ashing.

2.3 Fracture toughness test Compact tension specimens 3 5 m m X 3 6 m m were machined in both parallel (L) and normal (T) modes with respect to the mould fill direction (MFD) from the two mouldings (see Figs 2 and 3). Initial cracks were produced by first machining a notch with an included angle of 30 ° and then sharpening it by tapping with a razor blade at the root of the notch to produce a pre-crack ( ~ 0.4 mm in length). Tests were

Fracture toughness and impact behaviour of polyamide 6,6 111

section X-Z bottom 4 surface ~14mm g~te

Gate Gate ~ surface ~

72nun x (MFD)

(i) (ii)

0 ~ 3.4ram

~__~Z

x (M~)

(iii) (iv) Fig. 1. Schematic details of (i) PTO and (ii) plaque moulding, and photographs of (iii) PTO and (iv) plaque.

Table 1. Materials investigated

ICI grade Description Designation

Verton RFT0010EM(CB)

Verton RF70012EM(CB)

Maranyl A690

Maranyl A100

50% w/w long-glass-fibre-reinforced PA66, pigmented with 0-2% carbon 50LFPA black (CB), 10 mm long aligned fibres with a mean fibre diameter of 17/xm

60% w/w long-glass-fibre-reinforced PA66, pigmented with 0-2% CB, 10 mm 60LFPA long aligned fibres with a mean fibre diameter of 17/zm

50% w/w short-glass-fibre-reinforced PA66, pigmented with 0.2% CB, 50SFPA 0.2 mm randomly distributed fibres with a mean fibre diameter of 10/zm

PA Unfilled polyamide 66


<

~ e

T .

t ~

_t_

T

'- 315 -'

i

(i) (ii)

Fig. 2. Position of (i) transverse and (ii) longitudinal compact tension (CT) specimens in end-gated plaque

(drawn to scale and all dimensions in mm).

conducted on an Instron 6025 machine under ambient conditions of 25°C and 60% relative humidity, at a cross-head speed of 10mmmin -1. At least four specimens in both parallel and normal directions with respect to MFD were tested. Figure 4 shows the testing set-up.

Fracture toughness values were determined using ASTM E399:

P m a x f ( a ~ KQ=~W--~ \~/ where

(2+ a) (1 - t ~ ) 3/2 (0"886 + 4.64a - 13-32a 2 + 14.72a 3 - 5-6a 4)

and Pmax is the maximum load identified from the force/displacement plot (see Fig. 5), a = a/W, and a, B and W are as defined in Fig. 4(b). KQ is described as a provisional fracture toughness value, which becomes a critical value, KIc, if validation using several criteria as outlined in ASTM E399 in relation to specimen dimensions and forces are satisfied.

Forge

Test spec imen

Machined notch

/

Razor pre-craek

FOl'¢~ (a)

B = 3.4 - 4 .2

35

~ W = 24

E =,2 "

/ 16

. ( ) 1 , ~ 6 "

i T M i t . - q

36

Co)

Fig. 4. Schematic layout for testing of compact tension specimens.

<

(i) (ii)

Fig. 3. Location of (i) transverse and (ii) longitudinal CT speomens in centre-gated PTO (drawn to scale and all dimensions in mm).


0

IMPACT LOCATIONS

0.5 1 Fig. 5. Typical force/displacement plots for: (i) 50LFPA,

0.5

• ) 0 • 2 5

I I o 2 4 8

Displacemerl (ram)

(ii) 60LFPA, (iii) 50SFPA, and (iv) unfilled PA66.

2.4 Drop-weight impact test An in-house built instrumented falling weight impact apparatus described fully elsewhere 1° was fitted with a 10kg mass and a hemispherical nose of 12.7mm diameter. The impactor was dropped from a height to produce an incident impact velocity of 3 .8ms -1. Figure 6 illustrates the specimen support and impact locations. Each PTO moulding was clamped using seven bolts and a torque of 3.5 Nm (as employed in service), and subjected to impact at three separate positions. Graphs of force versus deflection were plotted (see Fig. 7) in real time.

3 RESULTS AND DISCUSSION

3.1 Microstructure A five-layer microstructural arrangement exists through-the-thickness of the mouldings, i.e. skin/shell(or shear)/core/shell(or shear)/skin (see Fig. 8). The skin and shell layers are not always distinguishable and, thus, an appearance of surface/core/surface layers may be more obvious. In general the in-plane fibre angle, oz, was 350-45 ° in the skin, 25°-35 ° in the shell layer (where high shear stresses exist) and 50o-60 ° in the core, very seldom did the fibre orientation in the core become completely transverse or at 90 ° to the melt fill direction, as is generally anticipated• Overall, a near random fibre orientation (o~ ~ 45 °) is present in the centre-gated PTO, while a higher fibre alignment parallel to the MFD exists in the end-gated plaque• Out-of-plane fibre angles do not vary as significantly as in-plane angles, typically from 10 ° to 19 ° across the moulding thickness, fibres are therefore considered to be planar in orientation, i.e. confined to the plane of moulding (x-y plane).

Table 2 displays the number average fibre length,

In = (E nili)/~, ni, for PTO and plaque mouldings (the values represent measurements taken from various regions of the mouldings). These data indicate that fibre breakage is more severe within the fiat plaque moulding, because of reduced nozzle and gating dimensions, and lower mould thickness. Increased fibre loadings from 50 to 60% (w/w) caused a reduction in the subsequent fibre length by 15 and 30% (depending on the mould type), due to the increased fibre-fibre and fibre-mould wall contact where the higher volume of fibres are forced through confined channels.

It should also be indicated that there was some variation in both the fibre length and the fibre content with position in the moulds and through the thickness

Fig. 6. Drop-weight impact test configuration.


4

3

0

(i)

0 4 8 Deflection (ram)

12 16

Fig. 7. Force versus deflection curves: (i) 50LFPA and (ii) unfilled PA66 (arrow indicates crack initiation).

Fig. 8. Scanning electron micrograph indicating skin, shell and core layers (50LFPA).

Table 2. Number average fibre length in both mouldings (combination of surface and core

measurements)

Number average fibre length, mm

PTO Plaque

Material

50LFPA 0.91 0.60 60LFPA 0.76 0.43 50SFPA 0-20 0.20

of the mouldings. The variation in fibre length was negligible in the mouldings containing the short-fibre grade and long-fibre grades which had suffered more fibre attrition, i.e. the plaque mouldings. In the long-fibre containing PTO mouldings, the average fibre length was greater in the core region by approximately 15%--a trend consistent with previous work.X1,12

Variations were also observed in the extreme regions of the mouldings compared with the rest of the mouldings (note that these were outside the areas representing the CT specimens). Most obvious were reductions of approximately 15% in the average fibre length and 7% in the average fibre volume content in the vicinity of the PTO gate and an accumulation of 10% by volume extra fibres near the back wall of the plaque mould for the long-fibre grades.

3.2 Fracture toughness Fracture toughness values (see Tables 3 and 4) for the short-fibre material do not vary significantly with the mould geometries, while for the long-fibre materials,

Table 3. Fracture toughness values for both transverse and longitudinal specimens from the plaque mouldings (standard deviations are given in parentheses)

Material K e (MN m -3/2)

Transverse Longitudinal

A (44°) a B (58°)" A (46°) a B (32°) a

50LFPA 7.36 (0.16) 9.82 (0-24) 7.79 (0-06) 6.36 (0.59) 50SFPA 8.24 (0.28) 8.66 (0.63) 8.63 (0.12) 8.18 (0.46) 60LFPA 8.13 (0.39) 9.01 (0.47) 8.79 (0.29) 7.18 (0.85) PA 4.30 (0.36) 4.24 (0.31) 4.27 (0.71) 4.63 (0.47)

a Average in-plane fibre orientation, a, values at crack tips.


Table 4. Fracture toughness values for both transverse and longitudinal specimens from the PTO mouldings (standard deviations are given in parentheses)

Material Ko (MN m -3a)

Transverse Longitudinal

A (44°) a B (46°) '~ A (46°) a B (44°) a

50LFPA 8.58 (0.12) 9.28 (0.45) 10.06 (0.28) 10.07 (0.51) 50SFPA 9.13 (0.21) 9.39 (0.10) 9.02 (0.19) 8.60 (0.08) 60LFA 10.09 (0.48) 9.81 (0.10) 11.60 (0.62) 10.65 (0.36) PA 4.21 (0.19) 4.02 (0.15) 4-04 (0.03) 4.45 (0-04)

a Average in-plane fibre orientation, a, values

(i) (ii)

(iii) (iv)

(v) Fig. 9. Fracture surface ahead of crack initiation region for Blongitud~a~ specimens: (i) 50LFPA PTO, (ii) 50LFPA plaque, (iii) 60LFPA PTO, (iv) 60LFPA plaque, and (v) 50SFPA

PTO.

at crack tips.

the specimens from the PTO yielded higher KQ values than those from the plaques--a reflection of the retention of relatively longer fibres in the PTO mouldings. The average fibre length, ln, for long-fibre materials is 34-43% less in the plaque than in the PTO. The corresponding reduction in fracture toughness is 15-25% (the greatest reduction both in In and K o is with the 60% w/w long-fibre material). This reduction in fracture toughness may be explained by examining a combination of effects: fibre length, fibre bundling and average in-plane fibre angle ahead of the crack-front (as dictated by the magnitudes of the skin, shell and core layers).

An indication of the longer fibres retained in the PTO can be seen in Fig. 9. Greater fibre and/or fibre bundle pull-out lengths are visible on the fracture surfaces of the PTO samples. This is clearly shown for both 50 and 60% w/w long-fibre material, where pull-out lengths greater than 3 mm are visible (see Fig. 9(i) and (iii)), compared with i mm pull-out lengths for samples representing plaque mouldings (see Fig. 9(ii) and (iv)). For comparison the equivalent short-fibre compound is displayed, with fibre lengths of approximately 100/zm visible on the fracture surface (see Fig. 9(v)). Recent work 13 has shown that composites containing fibre bundles exhibit better fracture toughness than the discrete fibre alternatives. Fibre bundles behave as single large fibres with effectively greater values of lc which should promote extensive fibre pull-out and, hence, contribute to fracture toughness.

A comparison of the fracture behaviour of all the materials used can be made by reference to Figs 5 and 10. Unfilled polyamide has a linear force/displacement behaviour, by contrast the fibre containing grades show non-linearity in the initial part of the curve and also an extension of the curve beyond the maximum load, indicating a progressive rather than a sudden failure. Figure 10 presents overall average KQ values, representing specimens from different types of mouldings and positions. Clearly, the incorporation of fibres improve the fracture toughness of the material considerably: average KQ values of approximately 4 M N m -3/2 and 9 M N m -3/2 are indicated for the polyamide and the fibre-reinforced polyamides,


12"

10"

8 ,,am ' G

4 -

2-

0 PA 50SFPA 5 0 L F P A 60LFPA

Mater ia l Type

Fig. 10. Fracture toughness, Ko, versus material type.

respectively. A greater improvement in the property is indicated in the higher fibre content grade (compare 50LFPA with 60LFPA).

There is also a significant scatter in the data for the fibre-reinforced grades, which may be related to the variations in the fibre length and orientation. A plot of the fracture toughness against the fibre orientation with respect to the crack direction is presented in Fig. 11, using the data for the plaque moulding from Table 3. Clearly the fracture toughness of the long-fibre grades shows a greater sensitivity to the fibre orientation at the crack tip.

An illustration of the influence of the fibre length on the fracture toughness, for a given fibre orientation, is presented in Fig. 12 by plotting the data representing the PTO and plaque specimens with an average in-plane fibre angle of 44 °. A projection of the 50LFPA data to unfilled PA indicates a K o value of 5 .4MNm -3a at 0.2mm fibre length, which is significantly below the measured value of K o ( --~ 8.7 MN m -3/2) for the Maranyl grade (i.e. 50SFPA).

12

10

8

4

2

0

X

O

I I ,0 , ; 40

aO

Fig. 11. Fracture toughness, KQ, versus average in-plane angle, a, between the fibres and the crack front for plaque mouldings: ( x ) 50SFPA, ( I ) 50LFPA and (O) 60LFPA.

12-

10"

8

4.

t 2 unfilled PA

0

O

0 012 0; 016 0; I .

Fig. 12. Fracture toughness, Ko, at a = 44 ° versus average fibre length, l,, for: ( I ) 50LFPA and (O) 60LFPA, ( x ) 50SFPA and ( t ) prediction for 0.2 mm fibre length using

measured data for 50LFPA and PA.

The superior performance of the Maranyl grade has to be attributed to the differences between the feedstocks: nominal glass-fibre diameter is 17/zm for Verton and 10/zm for Maranyl. Verton is pultruded, whereas Maranyl is extrusion compounded (enabling a more thorough and homogeneous mixing of the fibre and the matrix). Furthermore an examination of the micrographs in Fig. 9 shows that the fibres, pulled out of the matrix, are considerably cleaner in the Verton products--an indication of a difference in the strength of the fibre/matrix interface.

6

~ 4

~ 0

~ l ~ e region

~, Impact Locstion

Io +

Fig. 13. Average crack initiation force values for: ( I ) 50LFPA, (*) 50SFPA, (O) 60LFPA, and ( + ) PA. For

clarity some points have been offset.

I 0

0

OIt IO

÷ centre

Impact Location

Fig. 14. Average total energy values for: ( I ) 50LFPA, (*) 50SFPA, (O) 60LFPA, and (+) PA. For clarity some points

have been offset.

Fracture toughness and impact behaviour of polyamide 6,6

Table 5. Average impact crack initiation force and energy values and total energy values in PTO for filled and unfilled PA66

117

Property 50LFPA 50SFPA 60LFPA PA

Area A Area B Area A Area B Area A Area B Area A Area B

Initiation force, kN 2.4 2.3 2.8 2.6 2-5 2-5 1.9 1.7 Initiation energy, J 3.4 2-4 4.9 4.1 3.4 4.0 5.0 3.0 Total energy, J 21.6 20.0 21.0 17.8 21.1 20.7 5.0 3.0

3.3 Drop-weight impact behaviour Figure 7 shows typical force/deflection traces for filled and unfilled PA66. Impact properties in terms of the crack initiation force and the total energy (arising from initiation and propagation), are presented in Figs 13 and 14 and in Table 5. Crack initiation force values are 25-37% greater for the filled grades than unfilled PA66, (compare with the variation observed from fracture toughness measurements in excess of 100%), while the total energy fracture values are 300-600% greater for the filled compounds than unfilled PA66. There are negligible differences between long- and short-fibre materials and increases of long-fibre content from 50 to 60% w/w.

Similar to earlier fracture toughness measurements, failure of the unfilled specimen is in a brittle manner with long cracks propagating through the component causing it to shatter into several pieces. Filled materials tended to fracture as a laminated panel, where the damaged area of the lower face of the specimen was considerably larger than the directly hit upper face and showed the existence of internal fracture propagation between the layers. In comparison, the damaged area was more confined in the long-fibre grade mouldings with a typical appearance as shown in Fig. 15.

Total fracture energy and crack initiation force values are invariant to position within the moulding, with the exception of the centre region of the PTO, where although significant defects have been found, 14

greater crack initiation force and total energy values are produced due to the additional stiffening effects of the central rib, which is also the primary cause of the defects in the form of voids.

4 CONCLUSIONS

Fracture toughness values were determined using compact tension samples. A two-fold improvement in fracture toughness values was effected with the addition of fibres. Marginal improvements in K o values were produced with the introduction of longer fibres and with an increase in fibre content.

Values of K o were dictated, as would be expected, by the average in-plane fibre angles ahead of the crack front and the fibre length. Fracture toughness values were 15-25% higher in the PTO cover samples compared with the plaque samples for the long-fibre grades. The corresponding microstructures showed greater average fibre length (by 34-43%) and more fibre bundling in the PTO mouldings.

Under impact, brittle fracture and specimen shattering was observed with the unfilled polyamide 6,6 matrix. Short-fibre and long-fibre materials suffered localised rupture, the damaged area being smaller for the long-fibre materials. Impact crack initiation force and total energy values were increased by 25-37% and 300-600% by the inclusion of glass fibres, respectively. Crack initiation force and total energy values indicated negligible differences between long- and short-fibre materials and with increases of long-fibre content from 50 to 60% w/w.

50 mm

Fig. 15. Drop-weight fracture area (showing lower face) for 50% w/w long-fibre grade PTO.

R E F E R E N C E S

1. Blumentritt, B. F., Vu, B. T. & Cooper, S. L., Fracture in oriented short fibre reinforced thermoplastics. Composites, May (1975) 105.

2. Mandell, J. F., Huang, D. D. & McGarry, F. J., Crack propagation modes in injection moulded fibre reinforced thermoplastics. In Short Fibre Reinforced Composite Materials, ASTM STP 772, ed. B. A. Sanders. American Society for Testing and Materials, Philadelphia, PA, 1980, p. 3.

3. Jang, B. Z. & Lieu, Y. K., Fracture behaviour of short fibre reinforced thermoplastics: I. Crack propagation mode and fracture toughness. J. AppL Polym. Sci. (1985) 3923.


4. Lang, R. W., Manson J. A. & Hertzberg, R. W., Mechanisms of fatigue failure in short glass fibre reinforced polymers. J. Mater. Sci., 22 (1987) 4015.

5. Sato, N., Kurauchi, T., Sato, S. & Kamigaito, O., Mechanism of fracture of short glass fibre reinforced polyamide thermoplastic. J. Mater. Sci., 19 (1984) 1145.

6. Sato, N., Kurauchi, T., Sato, S. & Kamigaito, O., Microfailure behaviour of randomly dispersed short fibre reinforced thermoplastic composites obtained by direct SEM observation. J. Mater. Sci., 26 (1991) 3891.

7. Carling, M. J. & Williams, J. G., Fibre length distribution effects on the fracture of short fibre composites. Polym. Comp., 11 (1990) 307.

8. Karger-Kocsis, J. & Friedrich, K., Fracture behaviour of injection moulded short and long glass fibre Polyamide 66 composites. Comp. Sci. Technol., 32 (1988) 293.

9. Leach, D. C. & Moore, D. R., Failure and fracture of

short glass fibre reinforced nylon composites. Composites, 16 (1985) 113.

10. Barkley, D. & Akay, M., The design and evaluation of an instrumented impact tester. Polymer Testing, 11 (1992) 249.

11. Truckenmuller, A. & Abed, H. G., Injection moulding of long fibre reinforced thermoplastics: A comparison of extruded and pultruded materials with direct addition of roving strands. Poly. Eng. Sci., 31 (1991) 1316.

12. Bailey, R. S. & Rzepka, B., Fibre orientation mechanisms for injection mouldings of long fibre composites. Int. Polym. Process., 6 (1991) 35.

13. Kim, J. K. & Mai, Y. W., Fracture of CFRP containing impregnated fibre bundles. Comp. Sci. Technol., 49 (1993) 51.

14. Akay, M. & O'Regan, D. F., Generation of voids in fibre reinforced thermoplastic injection mouldings. Plast. Rubb. Comp. Process. Appl., 24 (1995) 97.

Documents

Fracture toughness and impact behaviour of glass-fibre-reinforced polyamide 6,6 injection mouldings