Failure and fracture of short glass fibre-reinforced nylon composites D.C. LEACH and D.R. MOORE (ICI Petrochemicals 0 Plastics Division, UK)

The failure and fracture behaviour of a range of reinforced nylon compounds have been studied using a number of different techniques. Toughness was measured by an instrumented falling weight impact method; fracture toughness was monitored by measurements of yield stress and a critical value for stress field intensity factor; deformational mechanisms were monitored using a volume strain/axial strain function together with the usual tensile stress/strain plot. Materials used in this work differed in type of matrix (nylon 6, nylon 6.6 and a toughened nylon 6.6), type of reinforcement (glass fibres mainly, but also spheres), fibre content (which was varied in the range 0 -50 weight %) and type of interfacial bond between fibre and matrix. It is apparent that whilst an individual technique might be adequate for monitoring performance, a combination of methods are needed to provide an understanding.

Key words: composite materials; fracture behaviour," injection moulding; thermoplastic resins; glass fibres

The use of short fibre-reinforced thermoplastics compounds in engineering applications is both growing and diversifying. Injection moulding components are emerging from motor industry applications to engineering tooling, where a cost effective set of properties are motivating advance. Two particular factors account for these trends: firstly, the prospects of a wider choice especially in the matrix materials; secondly, a greater awareness and understanding of material performance. It seems likely that further growth may be proportional to improved understanding, particularly where creative design of a compound and formulation optimization lead the way to satisfactory product performance. Our aims in this paper are therefore to discuss certain aspects of mechanical property evaluation which aid understanding in the assessment of new products.

In any study of injection moulding materials based on short fibre-reinforced thermoplastics, it is important to recognize the complications attendant on fibre orientation both within the plane and through the thickness of a moulding. Folkes' review 1 outlines many such considerations and provides a route to much of the detailed work. Dunn and Turner a recognized that the direction of fibre orientation in a moulding could not be simply represented and that due consideration needs to be given to the choice of test specimen and

moulding By the start of the 1980's, a clear view had emerged on evaluating stiffness and anisotropy of short fibre composites, but the assessment of toughness was far from well defined.

In order to provide an intrinsic measure of toughness, fracture mechanics techniques have emerged as the most constructive approach. Several workers have used such techniques in studies of short fibre-reinforced systems; for example Friedrich 3 in relation to fibre- reinforced polyesters, Mandell 4 on short glass and carbon fibre-reinforced nylon, polycarbonate, polysulphone and polyphenylene sulphide and Cann ~ in studies on GRP. It is well established that a single fracture mechanics parameter such as fracture toughness (Kc) or strain energy release rate (GO can often be inadequate in describing material toughness. Often, one of these parameters is supplemented with the yield stress (ay) because the size of a plastic zone (rp) describes more fully a propensity to crack initiation and growth:

rp °c (Kitty) 2 (1)

The exact relationship between plastic zone size and the material properties Kc and yield stress depends on the mechanism of fracture as well as other considerations. Nevertheless, a broad indication of the

0010-4361/85/020113-08 $03.00 © 1985 Butterworth f~ Co (Publishers) Ltd

COMPOSITES. VOLUME 16. NO 2. APRIL 1985 113

importance of these two fracture parameters has been identified 6 in the definition of ductility expressed as follows:

Ductility factor (Kc/o'y) 2 (2)

Toughness, in all its manifestations, can be evaluated by measuring the energy absorbed prior to failure (eg crack initiation) and fracture (complete crack propagation). When such measurements are conducted on various plate-type specimens, without introducing machined notches, then the failure/fracture process is influenced by the fabrication method employed in producing an injection moulding. The anisotropy causes heterogeneity through the mould thickness and dictates that a crack will grow in the weakest direction. Traditional impact methods such as the Charpy and lzod tests have only limited usefulness due to the dominance of the fibre orientation in the specimen on the measured toughness. Instrumented falling weight i m p a c t (IFWI) 7 has emerged as a satisfactory approach for studying toughness in this context.

An insight of the deformational mechanisms is also necessary if understanding is a key aim of a failure/ fracture study. Several techniques have featured in this role including volume strain measurement s , acoustic emission 9'~° and microscopy. In this work. use is made of a volume strain measurement technique because it can be simply conducted and adequately detects various mechanisms.

The subjects of this work are various nylon compounds. These include variations in the matrix material as well as variations in type and concentration of the reinforcing system. Experimental samples have been selected for illustrative purposes, rather than to provide a specific comprehensive evaluation. Our purpose is to demonstrate the potential in combining techniques for an evaluation of short fibre composites and to incorporate an ability to understand weaknesses and strengths of a system so that development of a product fit-for-purpose can be progressed.

EXPERIMENTAL DETAILS

Materials and mouldings

The materials used were all nylon compounds and included commercial grades and specially prepared samples. The matrices included nylon 6, nylon 6.6 and toughened nylon 6.6, unreinforced and reinforced with 10-50% by weight (w/w) short glass fibres. Two independently prepared samples were made of one material (nylon 6.6 + 33% w/w fibre). A sample of nylon 6.6 reinforced with 33% w/w glass spheres was also evaluated for comparative purposes. All materials were extrusion compounded and then injection moulded into ASTM bars and double-feed plaque mouldings 160 X 160 X 3.2 ram. Fig. I shows a sketch of the mouldings and nomenclature used.

The injection mouldings of fibre-reinforced nylon will be anisotropic and inhomogeneous at all levels of fibre content. In addition, the alignment of the fibres will not be simple. For example, fibres near the outer surfaces will be aligned in the direction of propagation of the melt front (due to a converging stress field during injection moulding), while fibres near the centre of the thickness will be aligned at right angles to the direction of propagation of the melt front (due to a diverging stress field during moulding). This skin/core

0 o

El

Injection points Injection point

? , Weld lin Line along which

polished section was taken b

0 o

Fig. 1 Injection mouldings used: (a) double-feed plaque, 160 x 160 mm, 3.2 mm thick; (b) ASTM bar, 215 mm long, 3.2 mm thick

model has been discussed by a number of workers and is reviewed by Folkes? The overall fibre alignment depends upon the ratio of skin to core, and this in turn is influenced by the thickness of the moulding and its geometry. In order to quantify this skin/core effect, a section was taken from one of the double-feed plaque mouldings of nylon 6.6 + 33% w/w glass fibre, along the line shown in Fig. 1. After polishing one face, the section was mounted on glass and the other face polished, and the resulting thin section examined in an optical microscope using transmitted light. Micrographs from various positions along this section are shown in Fig, 2. The micrographs show that over the majority of the section the skin/core proportion and fibre alignment are uniform, only being disturbed near the injection point and towards the end of the moulding. It can also be seen that the skin is strongly dominant with the core only occupying about 20% of the thickness. The same mould geometry and thickness were used for all the samples so we assume that skin dominates the fibre orientation. Consequently we can classify a specimen cut at 0 ° or 90 ° to the direction of propagation of the melt flow as, in broad terms, along or across the predominant fibre direction. Although this is a simplification, our purpose was to compare failure properties of different materials using a common moulding geometry while not ignoring the complex anisotropy and heterogeneity.

Mechanical tests

Failure and fracture studies were conducted using a number of mechanical properly techniques. The principal measurements included impact energy, fracture toughness, yield strength, modulus and volume strain. In recognizing the importance of moisture uptake on the properties of injection mouldings, all samples were stored dry prior to testing. All mechanical tests were run at 23°C in a temperature-controlled laboratory.

Thoughness was assessed in the IFWI method through measurement of the energy to initiate a crack and that to produce complete fracture. Full details of this technique are documented elsewhere. 7 In these experiments a ring support of 50 mm internal diameter was used, with a hemispherically nosed striker of

114 COMPOSITES. APRIL 1985

In ject ion point

Injection poirtt

a

160

I II !1

"".%% ; "

"%% '" ..'"'"'"

" "".. ' "-. "-.. Core , : ....'" ... y.." j.

1

Fibre orientation and distdbufion uniform over central section (120 ram)

- - - 2 o - - : . I II 1 1,~.2

~Core ~

'Skin'.~'

Fig. 2 Polished sect ion f rom double- feed plaque of nylon 6 .6 + 3 3 % w / w glass fibre. All d imensions in mm

C O M P O S I T E S . A P R I L 1 9 8 5 1 1 5

12.5 mm diameter and impact speeds of 5 m s -I . Test specimens were 80 × 80 mm squares cut from the four quadrants of the double-feed plaques.

Tensile modulus and volume strain were calculated from the tensile deformation of ASTM type bars on a universal testing apparatus (Instron) which incorporates strain measurements in three mutually perpendicular directions. The various types of extensometers for measuring axial strain (eA) and strain in the width and thickness directions (ie lateral strains e~v and et, respectively) are described in Reference 11. Volume strain (ev) is then determined from the expressions:

e v = {(1 + eA)(l + ew)(l + e t ) - l} (3)

These tensile deformations were conducted on ASTM bar specimens and hence the applied stress aligned with the predominant fibre direction. A test speed of 5 mm rain -~ was used for the tests.

The yield strengths of the samples were obtained in a plane strain compression test, following the procedure described by Williams) z This involved cutting strips of material in the 90 ° direction from the double-feed plaques and compressing them between dies of various breadths in order to compute a value of yield strength that would not be influenced by geometric or frictional considerations. A more usual tensile test could not be used because failure did not occur by net section yielding in the reinforced compounds; failure being accompanied by other mechanisms such as debonding and fibre pull-out

Typically five specimens were tested at four different die breadths using a cross-head speed of 0.5 mm min -1. Force/displacement plots enabled a peak force to be identified with a matrix yielding process and hence a yield stress to be calculated, but such stresses are highly dependent on die geometry. Consequently a plot of apparent yield stress vs die breadth (as described in Reference 12) enabled a stress at infinite die breadth to be extrapolated and thus a plane-strain yield stress was calculated.

Fracture toughness (Kc) was obtained through the analysis of single-edge notched (SEN) tensile tests, where the critical value of stress field intensity factor (fracture toughness) is determined from the expressionS3:

K c = crf Y a ~/z (4)

where a r is a gross fracture stress, Y is the geometry factor, detailed in Reference 13, and a is the notch depth.

Specimens were cut from double-feed plaques; tensile bars were cut such that the applied force was at 90 ° to the principal direction of propagation of the melt front. This enabled a machined notch to be cut along the direction of propagation of the melt front, thus the cracks were propagated parallel to the dominant fibre al ignment A fly-cutter, with a tip radius of 10 ktm, mounted on a milling machine was used to prepare the notches. In order to determine Ko at least eight specimens with notch depths varying between 10-60% of the specimen width were used.

Notched tensile specimens were simply clamped in an Instron machine but the calculation of the geometry factor (Y) contained in Brown and Strawley 13 assumes

that the tensile force is uniformly distributed across the width of the specimen. This assumption is consistent with pin-loading (provided that the distance between the pins is not less than three times the specimen width). It is pointed out by Brown and Strawley ~J that calibrations for SEN specimens which are eccentrically loaded in tension can be derived by superposition of results for axial tension and pure bending. Consequently, use of a geometry factor (Y) for pin- loading will yield a plot of stress vs (Yx/-a) -~ that does not pass through the origin, but intercepts the y-axis at a positive value commensurate with the bending stress generated during the opening of a crack. This is observed in all our fracture mechanics experiments and therefore K c is the slope of the plot, after a least squares fit, but without bias for the line passing through the origin. Independent verification of this is obtained by fracture toughness measurements on isotropic polymethylmethacrylate conducted in three- line bending, clamped SEN tension and pin-joined SEN tension. In all three tests K c was calculated at a value of 1.6 + 0.1 MPa m 1/2 and where the clamped SEN tensile tests were of the character just described.

DISCUSSION OF RESULTS

Instrumented falling weight impact

IFWI of plaques of short fibre-reinforced nylon compounds provide complex force/deflection curves. Often three or four cracks initiate and propagate in the plaques. Analysis of the results is focussed on the force to initiate fracture, together with the energies absorbed by the specimen in the initiation of a crack and the attainment of complete fracture, The details of analysis, typical force/deflection curves for an initiation/ propagation failure process and the handling of the statistical analysis are provided in Reference 7.

Scatter of impact data is an important consideration in the evaluation of materials, particularly those containing short fibre reinforcement. For example, in IFW| tests on a batch of 50 mouldings of a 33% w/w nylon 6.6 compound, it was observed that the failure mechanism was identical on a macroscopic scale. Analysis of the energy to initiate a crack showed that the highest and lowest value differed by a factor of two and that the coefficient of variation was 18%; quite a large value for a sample with 50 'identical' specimens. Consequently, meaningful impact evaluations can only emerge when relatively large quantities of specimens are tested, unless there are large differences in toughness.

The fracture of these specimens also gave valuable information on the predominant fibre direction. Crack propagation would be expected to occur in the weakest direction, which matched the direction of the melt front That is, crack propagation and alignment of the fibres in the surface layers was in the same direction. These observations for the experiments on mouldings with 33% w/w glass and other tests on mouldings with different glass contents confirmed our views on fibre orientation.

The effect of glass fibre content on the energy to fracture a plaque and initiate a crack for nylon 6.6 matrix composites is shown in Fig, 3. The unreinforced matrix exhibits relatively high energy absorption, but this is reduced considerably as soon as small quantities of fibres are present At fibre contents above about 10% w / w , the crack initiation energy appears to be

116 COMPOSITES. APRIL 1985

15

-a

~o

"8 5 E

Fig. 3

13

cture

Initiation

I I I [ 0 I0 20 30 40 50

Glass content (%w/w)

Effect of glass content on impact energy for nylon 6.6. Impact energy at 31% is displaced from 33% for clarity

S 3O ~-

ughened

~ ~ 20 ~ nylon 6.6

.... s - ' ' " Nylon 6

0 I I I I I 0 IO 20 30 4 0 50

Gloss content (% w/w)

Fig. 4 Effect of glass content on fracture energy for different matrices. The fracture energy at 31% is displaced from 33% for clarity.

insensitive to fibre content, but the total energy to fracture rises to a plateau with higher fibre contents. It is therefore implicit that the energy to propagate a crack increases with increasing fibre content. Mandell et al 4 show that cracks will propagate in a fibre avoidance mode through the matrix, provided that fibres are adequately bonded to the matrix. Consequently, as the fibre content is increased, then the crack route will be more convoluted, or more fibres must be fractured. Either explanation will lead to greater energy for crack propagation.

The choice of matrix also influences both energy to initiate and propagate cracks. This can be illustrated for short fibre-reinforced composites based on nylon 6, nylon 6.6 and toughened nylon 6.6, all examined over a range of glass fibre concentrations. The trends of impact behaviour with glass content are similar to that just described for systems based on nylon 6.6, although the magnitude of the effect is strongly influenced by the matrix material. Fig. 4 illustrates total fracture energy (ie initiation and propagation) vs glass content. Unreinforced nylon 6 appears less tough than nylon 6.6, although composites based on these two matrices become similar as the glass content increases. The matrix based on a toughened nylon 6.6, however, exhibits significantly higher energy absorption. These observations mainly relate to a crack propagation process because there are only small differences between these composites for the energy to initiate a crack (the toughened nylon 6.6 matrix requiring the highest energy for initiation).

The bond between fibre and matrix also affects the impact performance. Additional compounds containing 33% w/w glass fibres in a nylon 6.6 matrix were used to illustrate this effect. One compound had an amino propyl silane coating on the glass in order to produce a good interfacial bond, whilst the other compound had

a propyl trimethoxy silane coating which gave a poor bond. Fig, 5 illustrates histograms of the data for energy to initiate a crack, plotting proportional frequencey vs initiation energy. A student t-test confirms that there is a highly significant difference between the two sets since the probability of this difference in the means arising by chance is less than 0.1%. The beneficial effect of a good bond between matrix and fibre for good impact performance is confirmed, a conclusion in accord with the work of Seller et aZ 14

Fracture toughness tests

Fracture toughness (Kc) and yield stress were measured on samples of nylon 6.6 matrix-based composites, for glass fibre contents up to 50% w / w , on specimens taken from the 90 ° direction from the plaques. Data are illustrated in Fig, 6. In the yield stress test, the specimens at the highest glass content (50% w/w) fractured rather than yielded, therefore an extrapolation of the curve to 175 MPa is used instead of the plotted datum point. The plot in Fig. 6 shows that K c rises with the addition of fibres until a maximum occurs at 35% w/w glass content. This corresponds to 19% by volume of fibres. A similar peak has been reported by Friedrich 3 and Beaumont et al. 15

Davidge and Green ~6 suggest that the formation of a crack in a composite requires a critical amount of stored elastic strain energy within the reinforcing particles and the surrounding matrix. Consequently, crack formation depends upon the volume of the reinforcement. If the stress fields from adjacent particles overlap, then cracks will initiate and, for this condition to be satisfied, then a critical interparticle spacing can be defined. Davidge and Green 16 show

0.7

0.6

~' 0 . 5 -

~ 0 . 4 -

~ 0 . 3 -

~ 0 . 2 -

a. 0 .1

0 _ _ _ _ _ _ _ L _ _ 0 I

Fig. 5 Effect of Sample size = 20

' Poor ' bond

' Good ' bond i - - -3

I I 1 I

r - - '~ I I I L - - J I :-',,

~ _ _ t _ _ _ _ _ _ _ l . _ _ _ _ _ _ ~ 2 3 4 5 6

Crock initiation energy (J )

fibre/matrix bond on crack initiation energy at 23°C.

200 - [

150 - 4

IO0 T E 3

5o ~ 2

0

0 0

I - -

I I I I IO 20 30 40 50

Glass content (% w/w)

Fig. 6 Effect of glass content on yield stress (¢y) and fracture toughness (Kc) for nylon 6.6 based composites. The yield stress value at 31% is displaced from 33% for clarity

COMPOSITES. APRIL 1985 117

7 5 0

=L 50(

"~ 25(: Q

0 I I I I I 0 I0 20 30 40 50

Gloss content ( % w/w)

Fig. 7 Effect of glass content on ductil i ty factor at 23°C for compounds based on nylon 6.6

1 5 0

0

,oo v

b"

5O

Nylon 6.6

~ o ~ ~ ~ ~ ~ ' ~ " "~N y lon 6

Toughened nylon 6.6

0 I I I [ I 0 I0 20 30 4.0 50

Gloss content (% w/w)

Fig. 8 Effect of glass content on yield stress (Gy) at 23°C for compounds based on various matrix materials. The yield stress values at 19% and 23% are displaced from 20% and that at 31% is displaced from 33% for clarity

A

T E

no :E 2 v

I Toughened Nylon 6 6

Nylon 6

0 I I I I I 10 20 30 4 0 50

Gloss content (% w/w )

Fig. 9 Effect of glass content on fracture toughness (Kc) at 23°C for compounds based on various matrix materials. The fracture toughness at 31% is displaced from 33% for clarity

The effect of different matrices is illustrated in Fig• 8 for yield stress w glass content and in Fig, 9 for fracture toughness vs glass content. Nylon 6.6 shows the highest yield stress, while the toughened nylon 6.6 has the smallest. This observation would seem to be independent of glass content. A lower yield stress will lead to high ductility (as indexed through ductility factor) and in part accounts for the comparative behaviour already discussed for impact performance.

The graph of fracture toughness w" glass content (Fig, 9) shows that K c for the unreinforced nylon 6 is higher than that for nylon 6.6. However, introduction of fibre reinforcement ( eg 33% w/w) shows that nylon 6.6 based compounds are superior. The compounds based on toughened nylon 6.6 show a peak in K c, but at a lower glass content than for nylon 6.6 (20% w/w ~f35% w/w). This observation is probably in accord with the Davidge and Green model 16 previously discussed, but only a qualitative story now emerges. The toughened nylon 6.6 based compounds have a three phase structure - - glass, matrix and particles of toughening agent. Consequently, the critical interparticle distance will be reduced and as such this also reduces the fibre volume for opt imum performance. We observe this reduction (11% by volume at peak) but cannot quantify it by calculation.

The combination of yield stress and fracture toughness is illustrated as ductility factor ws glass content in Fig. 10. The lower yield stress and higher fracture toughness of toughened nylon 6.6 compounds compared with nylon 6.6 based compounds in the glass content range 10-20% w/w illustrates a higher ductility factor. This is in accord with the impact toughness data of Fig. 3. There is a suggestion, however, that at glass contents above 30% compounds based on nylon 6.6 might show superior performance. This demonstrates that it is insufficient merely to produce a highly toughened matrix because improvements in the toughness of the compound need to emerge not only from a lower yield stress but also from other factors such as improved fracture toughness. It is apparent, however, that toughness improvements at low fibre contents can be achieved by matrix modification, although usually accompanied by a stiffness and strength penalty.

The effect of fibre/matrix bonding agent on fracture toughness was examined for the same compounds used in the impact study. The material with the good bond gave a significantly higher fracture toughness; 2.7 MPa m 1/2 compared with 1.8 MPa m 1/2. These

that this relates to an interparticle spacing of at least twice the particle diameter. For fibre-reinforced composites where crack propagation is parallel to the fibres, then the particle diameter may be taken as the fibre diameter (11 /~m) and the critical interparticle spacing will then be at least 22 ktm. This leads to a critical fibre volume of 20% which is close to the maximum that we observe at 19%.

The combination of yield stress and fracture toughness enables a ductility factor to be calculated and a plot of this vs glass content is illustrated in Fig. 7. A maximum is again observed, but now at a slightly lower glass content, 32-33% w/w, which is more in line with the peak in Fig 3 for total energy absorbed for fracture. This corresponds to a fibre volume content of 15%.

1500

I000

Toughened

5 0 C

Nylon 6 .6

0 I I I I I 0 I0 20 3 0 4 0 50

Gloss content ( % w/w)

Fig. 10 Effect of glass content on ductil ity factor at 23°C for compounds based on various matrix materials. The ductil ity factor at 31% is displaced from 33% for clarity

118 COMPOSITES. APRIL 1985

experiments were repeated on another pair of compounds prepared by a more severe compounding route, but this time with cracks propagated across the dominant fibre orientation. Although the K c values differed, the conclusion that a good bond leads to a higher toughness was again confirmed (with values of 5.1 MPa m '/z and 3.7 MPa m ~/2 for the 'good' and "poor' bonds, respectively).

Although it is desirable to achieve good interfacial bonding between fibre and matrix, Wambach e t a P 7

suggest that this might be accompanied by a greater notch sensitivity. Our results do not confirm this and Mandell e t a # provide an explanation. They illustrate that compounds with a poor interfacial bond can have cracks propagating along that interface and thus debonding the matrix and the fibre. Materials with a strong interfacial bond have crack propagation through the matrix and thus a high fracture toughness.

Volumetr ic strain tests

The measurement of volume strain during a tensile test provides an insight into the various deformational mechanisms, as well as incorporating the usual monitoring of a tensile stress/axial strain function. The shape of a volume strain vs axial strain graph enables a qualitative appraisal of the deformational mechanism, particularly delineating between volume increasing processes (eg voiding, crazing, cracking, debonding, e tc )

and a shear banding process (yielding) where deformation occurs without volume increase. It can often be expected with filled or reinforced plastics that a volume increasing process will occur. A volume strain measurement alone cannot identify the specific mechanism, although complementary observations readily aid identification. Nevertheless, even without specific identification, a comparative guide of volume strain/axial strain function can be helpful. A number of illustrative examples are now discussed for various types of nylon compounds.

Nylon can be reinforced with glass fibres or spheres. The enhancement of modulus will depend on which glass system is used, even at a common reinforcement concentration. This is illustrated in Fig. 11 for tensile stress/axial strain plots for nylon 6.6 unreinforced and with 33% w/w glass spheres and fibres. Modulus, at a specific strain, is highest for fibre reinforcement, least for no reintorcement and intermediate for inclusion of glass spheres. Such data were obtained for ASTM-type tensile bar mouldings (stored dry), where the predominant direction of fibre alignment is in the direction of the applied stress. The volume strain data

~" 200 i

== -$

g I-

IOO

No gloss

I I I O I 2 3

Tensile strain (%)

Fig. 11 Tensile stress/strain curves for nylon 6.6 based compounds. Dry ASTM bars, 23°C, cross-head speed 5 mm min -1

"~ 33 % Fibres

0 I I I 0 I 2 5

Tensile strain (%)

Fig. 12 Volume strain/axial strain plots for nylon 6.6 based compounds. Dry material, 23°C, cross-head speed 5 mm min -1

obtained from the same experiments are illustrated in Fig. 12. It can be seen that there is a more pronounced volume increasing mechanism for the compound containing glass spheres compared with that for the compound containing fibres. The lower modulus of nylon compared with that of glass ensures that most of the strain occurs within the nylon matrix. Consequently, it is likely that the volume increasing mechanism is a debonding process between glass and nylon, this process of debonding being more pronounced with spheres than fibres. The similarity in volume strain vs axial strain functions between fibre- reinlbrced and unreinforced compounds suggests that the extent of debonding in the fibre-reinforced compound is quite small, since no debonding can occur in the unfilled material.

A debonding mechanism is strongly related to the adhesion between glass and matrix, and this in turn can be influenced by the coating applied to a glass fibre prior to preparation of an injection moulding compound. Volume strain measurements conducted on injection moulded ASTM tensile bars from the 'good' and "poor" bonded compounds referred to in previous sections are illustrated in Fig. 13. The compound with a poor bond between matrix (nylon 6.6) and glass fibres clearly exhibits a stronger tendency for a volume increasing mechanism, likely to be debonding between glass and matrix. The stress vs axial strain plots show a difference between these two samples. Both materials exhibit a similar stiffness at very small strains, but the compound with the good interfacial bond maintains a relatively high stiffness as axial strain increases. This is an alternative manifestation of debonding but an interpretation of the shape of a

o I 0.8

>

0 .2

O 6 Poor ' ,

0.4-- bond J ~f'Good' bond

0 I I I I 2 3

Tensile stroin ( % )

Fig. 13 Volume strain/axial strain plots for compounds based on nylon 6.6. Dry material, 23°C, cross-head speed 5 mm rain -1

COMPOSITES. APRIL 1985 119

stress/strain curve and the magnitude of the stress/ strain ratio introduces a different set of ambiguities and assumptions. On balance, the volume strain plot provides a more direct and informative view of the deformational mechanisms.

CONCLUDING COMMENTS

A single evaluation procedure would seem to be inadequate for an evaluation of failure and fracture of nylon composites. A combinat ion of techniques is seen to be more useful when collating a number of conclusions from this worlc

Impact toughness, assessed through IFWI, provides good resolution of initiation and propagation mechanisms and their monitoring, but does not explain trends with glass fibre content. The fracture mechanics results thus complement the impact data.

Total impact energy to fracture falls as fibre content increases from zero to about 15% w/w. This is explicable in terms of a larger increase in yield stress than that for K c. Consequently, the ductility factor plot of Fig. 7 shows good agreement with the impact energy plot of Fig. 3.

The general agreement between ductility expressed through ductility factor and toughness expressed through an impact energy is encouraging. It is not expected that complete accord between these measurements can be possible because they both contain failure and fracture mechanisms manifesting themselves in different ways. Nevertheless, a full compar ison of the two techniques, as illustrated in Fig. 14 for all the materials, does provide a positive correlation.

Evaluation of an individual failure mechanism remains important. For example, the process of crack propagat ion as monitored through fracture toughness appears to agree with the Davidge and Green modeU 6 reinforcing the belief that an unders tanding of toughness is at hand.

Other deformational processes, such as debonding` can be identified through a volume strain v axial strain plot and this technique confirms aspects of material behaviour concerning reinforcement type and glass- fibre coating.

The various material changes incorporated into our samples, whether glass concentration, matrix type or interfacial bond, have demonstrated the value of combining various failure and fracture techniques. It is not enough, however, merely to moni tor a process: it is apparent that some complementary unders tanding is also important.

ACKNOWLEDGEMENTS

The authors acknowledge the constructive contributions from J.N. Gaitske]l in discussions and provision of sample mouldings; from M. Whale and J. Hastings in mechanical testing; from A.D. Curson in optical microscopy: and from Pilkington Bros plc in certain glass fibre samples.

This paper is based on one first presented at the Concerence 'Testing Evaluation and Quality Control o f Composites', held at the University of Surrey, UK. on 13-14 September 1983.

1500 -

::k

t o o o -

"~ 500 - O

Fig. 14

o < 15% fibres x 3 0 - 3 3 % fibres o 20% f ib res A 4 0 - 5 0 % f ib res

X i'1 0 0

0 0

XX x X

X A 0 X X A

O

I I I I I 5 10 15 20 25

F r o c t u r e energy (d)

Relationship between ductility factor and impact fracture energy fo r a range of compounds

REFERENCES 1 F o l k e s , M . 'Short Fibre Reinforced Thermoplastics' (Rsch

Studies Press, J Wiley & Sons Ltd, 1982) 2 Dunn, C.M.R. and Turner, S. 'The characterisation of

reinforced thermoplastics for industry and engineering uses' NPL Conf Proc April 1974

3 Friedrieh, K. "Microstructure and fracture mechanical properties of short fibre reinforced thermoplastics PET' PRI Conf Proc 'Deformation YieM & Fracture of Polymers] April 1982

4 MandeU, J.E., Darwich, A.Y. and MeGarry, F.J. "Fracture testing of injection-moulded glass and carbon fibre- reinforced thermoplastics' ASTM STP 734 edited by Chamis (American Society for Testing and Materials, 1981) pp 73-90

5 Cann, R.J. 'A fracture mechanics approach to static and fatigue failure in glass reinforced plastics' PhD thesis (Nottingham University, U~ 1977)

6 Moore, D.IL, Hooley, C.J. and Whale, M. "Ductility factors for thermoplastics" Plastics and Rubber Proc & Appl l No 2 (1981)

7 Gutteridge, P.A., Hooley, C.J., Moore, D.R., Turner, S. and Williams, M.J. 'A versatile system of impact tests and data' Kunststoffe 72 (1982) p 9

8 Bneknall, C.B., Clayton, D. and Keast, W. "Rubber- toughening of plastics" JMater Sci 8 (1973) p 514

9 Jackson, J.R. "Acoustic emission from short fibre GRP composites" Proc Conf on Interfaces in Comp Mater, Liverpool UK, April 1981 Welters, J. "Influence of shape and bonding of glass particles on the acoustic emission behaviour of polycarbonate' XI EWGAE Meeting Milart Italy. November 1982 Barrie, I.T., Moore, D.R. and Turner, S. "Developments in the generation and manipulation of mechanical properties data' Plastics and Rubber Proc & Appl 3 (1983) p 293 Williams, 3.G. "Plane strain compression testing of polymers' Trans J Plast lust 35 (1967) p 505 Brown, W.F. and Strawley, J.E. 'Plane strain crack toughness testing of high strength metallic materials' ASTM STP 410 (American Society for Testing and Materials, 1966) Seiler, E., Dorst, H.G. and Theysohn, IL "New developments in glass fibre reinforced thermoplastics' 5th Euro Plastics & Rub Conf ParL~ France, 1978 Beaumont, P.W.R., Jolliffe, V. and Gore, D. 'Fracture of particulate composites based on PMMA' Fracture 3 (1977) pp 1015-1023 Davidge, R.W. and Green, T.J. "The strength of two-phase ceramic/glass materials" J Mater Sci 3 (1968) p 629 Wambach, A., Trachte, K. and DiBenedetto, A. "Fracture properties of glass filled polyphenylene oxide composites" J Comp Mater 2 No 3 (1968) pp 266-283

10

11

12

13

14

15

16

17

A U T H O R S

The authors are with the Analytical and Polymer Science Group, Research and Technology Department, Imperial Chemical Industries plc, Petrochemicals and Plastics Division, PO Box 90, Wilton, Middlesbrough, Cleveland, TS6 8JE, UK. Inquiries should be directed to Mr Leach.

120 COMPOSITES. APRIL 1985