Thermoplastic impregnated fiber bundles: Manufacturing of laminates and fracture mechanics characterization

Composites Science and Technology 33 (1988) 97-120

Thermoplastic Impregnated Fiber Bundles: Manufacturing of Laminates and Fracture Mechanics Characterization

K. Friedrich Polymer and Composites Group, Technical University Hamburg-Harburg, 2100 Hamburg 90, FRG

T. Gogeva & S. Fakirov Laboratory on Structure and Properties of Polymers, University of Sofia, Bid. A. Ivanov 1, 1126 Sofia, Bulgaria

(Received 6 November 1987; accepted 18 February 1988)

ABSTRACT Linear flexure-response-studies, through-thickness fracture toughness tests ( Kc), and interlaminar mode I and mode H fracture energy measurements were carried out with different laminates of a carbon fiber/thermoplastic polyamide 12 composite system. Specimens were preparedfrom fiber bundles interspersed with polymer powder and surrounded by a polymer sheath. The results, which must be considered as preliminary data because of very limited availability of specimen material, reflect an overall good fracture toughness profile of the different laminates tested. This is also highlighted by thefracture surface micrographs achieved ~¢ith SEM-analysis.

NOTATION a aN B C E

Crack length (mm) Notched Charpy fracture energy (kJ m-2) Specimen thickness (mm) Compliance of DCB-specimens (mm N - 1) Elastic modulus (MPa)

97 Composites Science and Technology 0266-3538/88/$03'50 © 1988 Elsevier Science Publishers Ltd, England. Printed in Great Britain

98 EFlex

F Fcrit

Fmax

Gc GIc Gnc

h

Ko

/ L P

Vr W

W x

6 o O'max O'o.ma x O'fF

K. Friedrich, T. Gogeva, S. Fakirov

Flexural modulus (MPa) Load (N) Critical load (N) Maximum load (N) (Through thickness) fracture energy (kJ m-2) (Interlaminar) fracture energy (Mode I) (kJ m-2) (Interlaminar) fracture energy (Mode II) (kJ m-2) Thickness of beam (mm) (Through thickness) fracture toughness (MPaw/m) Length of the mold (mm) Span (mm) Pressure (MPa) Fiber volume fraction (%) Width of the mold (ram) Specimen width (mm) Displacement (mm) Deflection (mm) Diameter of fibers (mm) Maximum breaking stress (MPa) Strength of 0°-laminate (MPa) Fiber tensile strength (MPa) 1 INTRODUCTION

The trend to use thermoplastic polymers as matrices in high performance composites arose from several disadvantages of the thermosetting resins primarily used at present. Thermoplastics usually result in better residual compressive strength after impact than thermosetting resin composites

>= .=_E

z=£

Toughened Thermosets 0

Thermoplastic dard oxy

oughened BMI or PI

Hot/Wet Compressive Strength Fig. 1. Schematic illustration of residual compressive strength after impact vs hot/wet compressive strength of different thermosetting resin composites by comparison with thermoplastic matrix composites (after Ref. l).

Thermoplastic impregnatedfiber bundles

99

which have been specially designed for good hot/wet compressive strength. They also exhibit better hot/wet compressive strength properties than thermosets designed to retain high residual compressive strength after impact (Fig. 1). Thermoplastic matrices reinforced with continuous glass or carbon fibers also possess better resistance to interlaminar crack propagation. In the case of carbon fiber/polyetheretherketone (CF-PEEK) composites, for example, the values of mode I and mode II interlaminar fracture energy (G~c,Gnc) are about 10 times those of corresponding epoxy CF-PEEK (Mode II) CF-PEEK (Mode I ) ~ ~ i,. A

"E 1,1..,_. 0.1

T 0'0110-10

I 10-9

I 10 -8

I 10 -7

1 10-6

I 10-5

Crack Tip Displacement Rate (m/see) Fig. 2. Mode I and mode II interlaminar fracture energy data of carbonfibre/PEEK and carbonfibre/epoxy composites as a function of crack tip displacement rates (after Ref. 2).

matrix composites (CF-EP). This is summarized in Fig. 2 for data obtained over a wide range of crack tip displacement rates. 2 Additional advantages achieved with a thermoplastic in comparison to a thermosetting resin matrix are: I. 2. 3. 4.

thermoformability after consolidation 3'4 weldability s recyclability new processing techniques, e.g. laser consolidation during fiber placing 6

One of the problems with thermoplastics compared with thermosets in the past was their high melt viscosity which makes it rather difficult to impregnate fiber bundles so as to ensure that the fibers are well wetted by the polymer in the final prepreg or in the laminate. However, several techniques have been developed over the last few years which overcome this difficulty.7 -9 One method is the use of low-viscosity, thermoplastic polymer solutions to thoroughly impregnate the fibers (with the drawback that the solvent has subsequently to be removed). Other methods start with

100

K. Friedrich, T. Gogeva, S. Fakirov

intermediate forms of the thermoplastic polymer, e.g. as films or fibers, which on melting are converted into a matrix embedding the reinforcing fibers (Fig. 3(a) and (b)). Further alternatives are: (1) fiber bundles may be pultruded through an extruder die, ending up as stiffextruded polymer rods or strands with reinforcing fibers in between; (2) fiber bundles may be infiltrated with solid polymer powder and surrounded by a thin polymer sheath (Fig. 3(c) and (d)). The latter intermediate material form can be produced by guiding fiber bundles through a fluidized bed of polymer powder. Spreading the individual fibers in the bundle enables an uptake of the solid powder particles between the individual fibers. Subsequently, the powder-infiltrated bundle is coated with a thin polymer sheath by passing it through an extruder die. The latter technique was used by Atochem (France) for producing the intermediate material forms needed for the study reported here. With carbon fibers as reinforcement and polyamide-12 as the thermoplastic matrix we have investigated how laminates could be made from this intermediate material form (on a laboratory scale) and what fracture mechanics performance could be expected from such composites. Co- or Intermingled Fibers

Film Stacking

III" u

e ~

u

-

-

-

Polymer Fibers

Polymer

Films

(a)

(b)

Powder Impregnated Bundle

Melt Impregnated Bundle

Fiber Bundle

Reinforcing

Fibers ~J li':

Extruded'~~ Polymer Powder Polymer

Sheath

Poly_mer j H°a I I

ifl']

1 i I I

(c) (d) Fig. 3. Several intermediate material forms for high-performance thermoplastic matrix

composites.

Thermoplastic impregnatedfiber bundles

101

Fig. 4. Fiber/matrix arrangement in the intermediate form: (a) overview showing an end of a fiber bundle surrounded by the polymer sheath; (b) higher magnification of the individual fibers in the bundle (with polymer powder in between).

102

K. Friedrich, T. Gogeva, S. Fakirov

2 E X P E R I M E N T A L DETAILS

2.1 Material The material investigated is a commercial product of Atochem (France), i.e. FITt, consisting of 68% carbon fibers (Torayca 6 K) and 32% polyamide- 12 by weight. 1° This corresponds to a fiber volume fraction between 50 and 55%. The intermediate form in which fibers and matrix were combined was that of a roving (Fig. 4). The components were arranged in the following way: (i)

A continuous bundle of 6000 individual carbon fibers (e = 8/tm); with (ii) PA-12 powder between them (particle diameter similar to that of the individual fibers in order to achieve better infiltration); and (iii) a PA-12 sheath of thickness about 10/~m around the whole bundle.

Consolidation of this intermediate material form into prepreg sheets of thickness 0.3 to 0-4 mm was carried out in a hot press using a steel mold of

Pressure

Temperature~ M

,i \ "~'Steel Foils Fiber Bundles

Fig. 5.

Geometry of the mold for consolidating fiber/thermoplastic matrix bundles into prepregs and laminates.

dimensions l l 0 m m x 50mm (Fig. 5). For the manufacturing procedure, about 90 threads of unidirectionally oriented fiber bundles were placed (unfixed) into the mold, either parallel or perpendicular to the length axis (l) of the mold. After heating the system to 210°C a compressive load of 30 kN was applied over a period of 15 minutes (mold pressure p = 5-5 MPa). The consolidated sheet was then cooled under the same pressure down to room temperature within 5 to 6 minutes. Steel foils were placed between the mold plates and the sample in order to achieve easy release from the mold cavity and smooth specimen surfaces. t FIT" = fibre impregn6e thermoplastique.

Thermoplastic impregnatedfiber bundles

103

(a)

(b)

Fig. 6.

Polished section of the [90, O, 90, ff]~ laminate: (a) overview; (b) higher magnification showing that fibers are well surrounded by the polymer matrix.

104

K. Friedrich, T. Gogeva, S. Fakirov

In a following step three types of laminates were produced from these prepregs (using the same compression molding technique): (a)

[0], laminates with thickness of about 1.5 m m and 3 m m respectively (i.e. 4 or 8 plies) (b) 1-90, 0, 90, 0] s laminates, about 3 m m in thickness (Fig. 6) (c) [90]s laminates (thickness 3 mm)

2.2 Mechanical testing Because of the limited availability of test material and of specimen geometry, the different laminates could not be tested comprehensively, i.e. with respect to their total mechanical property profile. Only a few selected tests were carried out which were expected to provide a good feeling for the potential of this material group in the area of composites with a requirement of high fracture toughness. Fracture toughness (K~) studies were performed with compact tension (CT) specimens machined from the different laminates (thickness = 3 mm). The unidirectional ones were notched either parallel (L) or transverse (T) to the fiber direction (Fig. 7). All machining and notching were done with a

t~ 36 ~ u : (a)

Kmfli3~ m:m:~ 0'~i LQaJ=29"~t" _

~ (b)

~

A

B=3

-~

Fig. 7. (a) Compact tension specimens with different crack directions machined from the [0Is laminates; (b) CT specimen geometry.

diamond saw. Prior to testing, the notch was sharpened with a razor blade. Tensile loading of the CT-specimens was performed on a ZWICK 1445 type static testing machine at room temperature and a crosshead-speed of 1 mm m i n - 1. From the resulting load-displacement curves the critical load at crack instability, F , , , was used to calculate a fracture toughness value, Kc, according to the equation: Kc -

Fcrit

(l)

Thermoplastic impregnated fiber bundles

105

where B = specimen thickness (2.7-3.1 mm), W = specimen width (29 mm), a = actual crack length at the beginning of each test, and Y(a/W)=polynomial geometrical correction factor. Plaques with thicknesses of ,,~ 1.4mm were used to prepare double cantilever beam (DCB) and end-notched flexure (ENF) specimens. Both types are typically used for determining interlaminar fracture energies under mode I and mode II crack opening conditions, respectively. A Teflon film (65/~m thick) was placed between the plaques (through-width) which were then pressed as mentioned above. Figure 8 shows the different specimen geometries corresponding to the specific types of fracture energy tests. Test

/ht

~-a -~ [

I

I

~-a~ © Iq

I L

~'~i ~

L

Double (DCB) End Notched Flexure (ENF)

Cantilever Beam

Fig. 8. Geometry and loading arrangement of the DCB and ENF specimens used for the interlaminar fracture toughness studies.

conditions were essentially the same as for the K c test, i.e. room temperature and 1 m m m i n -1 crosshead speed. The mode I fracture energy (G~c) calculation followed the compliance method as described in Ref. 11: F2rtt ( d C ' ~

G,¢ = ~

\d~-a,./

(2)

where F . i , is the critical load, W is the specimen width, and dC/da is the change in specimen compliance with crack length, experimentally determined according to the recommendations given in Ref. 11. From the load vs displacement curves obtained from the E N F specimens, the critical strain energy release rate in mode II loading (mode II fracture energy, Gnc) was estimated from eqn (3): 9F2rttCa 2 Gn~ = 2 W(2L 3 + 3a 3)

(3)

where F~rit is the critical load, C is the compliance evaluated from the crosshead displacement, a is crack length, W is the specimen width, and L is

106

K. Friedrich, T. Gogeva, S. Fakirov

the span, i.e. the distance between the central loading pin and the outer supporting pins. Tests were carried out on a three-point flexure testing jig at a displacement rate of 0.5 mm min- 1. Prior to the mode II loading, a natural crack was created by wedge-opening the in-laminated crack, i.e. according to the mode I precracking technique described by Carlsson et al. 12 Laminar flexure response (LFR) was determined by a three-point flexure test at room temperature and a crosshead speed of i mm min-1. As the recommended span-to-thickness ratio 11 could not be obtained with the present specimen geometry, a certain contribution of interlaminar shear deformation affects the resulting flexure data. This effect is neglected however in the calculations that follow. The ultimate flexural stress was obtained from the maximum of the load-versus-strain plots using eqn (4): 3FmaxL O'max -- (2 Wh 2)

(4)

where FmaX is the maximum load, L is the total support span, h is the thickness of the beam (plate thickness, B) and Wis the width of the beam. In a similar way, the flexural modulus was calculated: FL 3 EF|ex -- 4 Wh3•

(5)

where L, W and h are as defined above, F is a load in the linear range of the load-deflection curve, and 6 is the deflection in the center of the specimen at this load (neglecting shear deformation11).

2.3 Fractography Studies on the fracture mechanisms were carried out by analyzing fracture surfaces of CT, DCB and ENF samples in a Leitz 166 T-scanning electron microscope (SEM). Prior to viewing, the surfaces were coated with a thin gold layer using a Leitz sputtering chamber.

3 RESULTS AND DISCUSSION

3.1 Laminar flexure response Results of flexure studies on the three different laminates are listed in Table 1. At a relatively small span-to-thickness ratio, L / B , of about 15, the values achieved with the two 0 ° specimens (1 and 2) and those achieved with the [90, 0, 90, 0]~ specimens (3 and 4) represent the reproducibility of the materials data, even from this rather primitive laminates manufacturing

Thermoplastic impregnatedfiber bundles

107

TABLE 1

Flexural Response Data of CF-PA12 Laminates No.

1 2 3 4 5 6

Fiber orientation

F=,,,,

[0]8 [O]a [90,0, 90, 0]s [90,0, 90, O]s [90,O,90, O]s [90] a

475 445 293 245 250 44

(N)

~m.,, EEl=. ( M P a ) (GPa)

669 627 332 322 410 58

L (mm)

W (mm)

B (mm)

Span to thickness ratio L/B

40 40 40 50 65 40

6-3 6'3 6'3 6'3 6"6 6"3

2"6 2-6 2.9 3'0 3"0 2'7

15"4 15"4 13"8 16"7 21"6 13-8

104 108 19.5 20"9 46-4 4"2

procedure. A comparison of the data achieved with the three different laminates at LIB = 15 yields a linear trend in the flexure strength values, whereas the flexural modulus shows a progressive trend (Fig. 9). Using unreinforced matrix and fiber data for a prediction o f the flexural modulus of the [90, 0, 90, 0Is laminate, according to the laminate theory concepts, ~3 gave, however, a higher value (Era, x -- 51"5 GPa) than the one experimentally determined. This is probably due to the span-to-thickness ratio dependence o f the modulus at lower LIB ratios, as described by Carlsson and Pipes. x i In the present case, one sample was studied at an LIB ratio o f a b o u t 22, and the measured modulus was clearly higher than those of the two others (cf. specimens 3 and 4 relative to 5). The recommended span-to-thickness ratio of at least 3211 for the test specimens used here was not achieved however (Fig. 10). Regarding the strength values of, for example, the [0], laminates, it

EFlex

(GPa)

100

60--

/

0~ ~

j

600

/ / ' /

Gm~x

40 ~ 20 - / " ~ [901e

(MPa) / - /

80 -

Fig. 9.

(~max

(L/B ,~ 15)

~

//Z~

/ / /

/

/ ~

/

~0 ~

400 -- 200

\E

1

0 [90,0,90,0] s

[0] 8

Flexural modulus and maximum flexure stress at break as a Function of laminate

structure (span-to-thickness ratio ~ 15).

108

K. Friedrich, T. Gogeva, S. Fakirov

theoretical prediction

50

f

40 EFlex (GPa)

30 20 10 0

. . . . . . . . . . . . . . . . .

recommended L/B

I

I

I

I

I

I

5

10

15

20

25

30

Ip 35

L/B

Fig. 10. Flexural modulus versus span-to-thickness ratio for the [-90,0, 90, 0]s laminates, along with the theoretical prediction and the recommended span-to-thickness ratio after Ref. 4. becomes obvious that these are clearly lower than those achieved by Carlsson et al. I t with [0] samples well prepared from C F - P E E K prepregs (commercially available from ICI, U K , as APC-2). On the basis of the same volume fraction (APC-2 has a b o u t 62% by volume carbon fibers), the O'max values of the F I T material are lower by a factor of 2. One explanation of such a difference in this mainly fiber-dependent property may again be found in the limited specimen geometry used here. But another reason is definitely the deviation of many of the fiber bundles in the prepregs from the ideal direction in which they were placed. This could be seen when breaking [0] specimens parallel to the fiber direction along interlaminar planes (Fig. 11).

Fig. 11. Interlaminar fracture surface, showing bundles ofmisoriented fibers(by an angle of 20°) in the [0]o laminates (arrow).

Thermoplastic impregnatedfiber bundles

109

3.2 Through-thickness fracture toughness Kc F r a c t u r e o f the C T samples perpendicular to the loading direction could only be achieved with the bi-directional laminate ([90, 0, 90, 0Is at a critical load, Fcr~t = 2002 N) and with the [90°Is laminate (identical to [-0]s laminate with starter notch parallel to the fibers, i.e. L-cracks at a critical load, Fcr~t= 385 N). On the other hand, T-cracked specimens b r o k e after shear d e f o r m a t i o n parallel to the fibers, starting from the root o f the razor notch in the direction o f the applied load (Fig. 12). This is the reason for the rather

........ ~

)%1 !

(a)

¸¸¸ ¸¸¸

(b)

(c) Fig. 12. Macroscopic fracture response of CT specimens with different laminate structure: (a) [90 °] laminate (identical with a [0°]8 laminate with notch in L direction, i.e. parallel to the fibers); (b) [90,0,90, 0]s laminate with notch parallel to the fibers in the 90°-plies; (c) [0"]8 laminate with notch in T direction, i.e. perpendicular to the fibers (arrow indicates crack deviation).

110

K. Friedrich, T. Gogeva, S. Fakirov

low load values measured for the [0] a laminates (Fort, = 525N). Higher magnifications of the fracture surfaces of the [90]8 and [90, 0, 90, 0]s specimens are represented in Figs 13 and 14. Those fibers which were oriented parallel to the main fracture plane look very clean, i.e. surrounding matrix material must have been detached from them very easily during breakdown of the composite. This also reveals rather inadequate wetting of the fibers by the thermoplastic matrix and by the processing technique used here. The micrographs give further indication of a high degree of matrix deformation prior to final failure of the samples.

Fig. 13.

CT specimen fracture surface of a [90°]8 laminate showing matrix deformation

between the fibers and poor fiber/matrix bonding.

The lower magnification SEM-micrograph of Fig. 14(a) demonstrates the total appearance of the fracture surface of the [90, 0, 90, 0]s laminate (crack direction from lower left to upper right; specimen edges upper left and lower right). Fibers in the 0 ° layers were fractured at some distance away from the major fracture plane of the 90 ° layers, a typical indication of rather poor bonding between the fibers and the polymer matrix. They were then pulled out of the fracture surfaces mostly in form of bundles rather than individually which can be seen in higher magnification on Fig. 14(b). K c values, calculated from the critical loads and the initial crack length were clearly a function of the laminate structure. The rather high value for the

Thermoplastic impregnatedfiber bundles

111

Fig. 14. (a) Total fracture surface of a [90, 0, 90, 0]~ laminate CT specimen at low magnification (b) section of laminate at higher magnification.

K. Friedrich, T. Gogeva, S. Fakirov

112

[90°]3 laminate ( ~ 5 MPa ~ / ~ ) reflects the capacity of the matrix between the fibers to absorb energy during intralaminar crack propagation. The value is, however, lower than that calculated for unfilled polyamide 12 from notched Charpy and tensile modulus data ( 7 - 8 M P a x / ~ ) . l-As This is typical for tough thermoplastics used as matrices for high performance composites, e.g. when comparing their fracture energy with the interlaminar fracture energy of their counterparts.l 6.1 v The reason for this is simply the fact that the damage zone size in front of the crack cannot be built up in the same way in the composites (because of the rigid fibers) as would usually be found in the unreinforced polymer. This is opposite to the trend observed in more brittle epoxy systems in which the fibers involved in crack bridging and other effects normally cause higher values for the composite by comparison with the unfilled matrices.1 s The addition of plies with fibers transverse to the intended crack direction (from [90°]3 to [90,0,90,0]3 ) leads to an enormous increase in crack resistance of the composites (Kc = 30 MPa x//-m). This value is in the same range as those measured for other carbon-fibre reinforced composites laminates with similar fiber arrangements? 9.zo The rather high value is especially based on the poor bonding between the fibers and the matrix which enables a significant contribution of the 0 ° layers to energy absorption by fiber fracture, fiber matrix debonding, and pull-out. Stronger bonding in these systems would result in lower K c data. 2~ The individual amounts of energy absorption by the different layers in the laminate can be estimated when converting the Kc values into G¢ values, according to the following relationship: Gc

=

(6)

KZ~/E

where for E the flexural moduli were used for simplicity (Ego = 4.2 GPa; E9o,o,9o,o = 50GPa). The results plotted in Fig. 15 versus the laminate structure (with extrapolation to a [0 °] laminate with crack in the transverse 40

35 t 30

(kj/~ 2) 20 I--

............................................................................... i .......

~

......................................................................

coo,,out,on

Ol [90] S

~3 C°ntributi°n 90°'layers [90,0,90,0] S [0] 8

Fig. 15. Fracture energy (G~) data, calculated from the K c and flexural modulus data of the [ 9 0 ] 8 and [90, 0, 90, 0]~ specimens and their extrapolation to a [0c]8 laminate structure ([90]8 ~ eight 90 ° layers, in comparison with [90, 0, 90, 0]~ with four 90 ° and four 0 ° layers).

Thermoplastic impregnatedfiber bundles

113

direction and assumed fracture transverse to the fibers). Each 0 ° layer absorbs about five times more energy during fracture of the composite than the 90 ° layers do. 3.3 Interlaminar fracture energies G~c and Gnc Figures 16(a) and (b) illustrate typical load--displacement responses from the DCB and E N F specimen geometries of the I-0°] a laminates. The initial linear part is followed by a distinct non-linear portion and subsequent stable (I) or unstable (II) crack propagation. The maximum of these curves was used for the calculation of the interlaminar fracture energy values. Under mode I F [N]

250

Mode I

200

150

100

50

t

2

3

4

5

X[mm]

(a) F [N]

500

Mode I I

400

300

200

100

i

0

I

~

I

3

X[mm]

(b) Fig. 16.

Typical load displacement curves obtained with interlaminar crack growth specimens under (a) mode I and (b) mode II conditions.

114

K. Friedrich, T. Gogeva, S. Fakirov

(a)

(b) Fig. 17. Low (a) and high (b) magnification micrographs of the fracture surfaces of DCB specimens (mode 1 failure).

Thermoplastic impregnatedfiber bundles

115

(a)

(b)

Fig. 18. Typical appearance of the fracture surfaces of EN F specimens failing under mode II (shear) conditions.

116

Fig. 19.

K. Friedrich, T. Gogeva, S. Fakirov

Comparison of (a) tensile and (b) shear deformation of the matrix, as found on the DCB and E N F specimens, respectively.

117

Thermoplastic impregnated fiber bundles TABLE 2 Mode II Fracture Energy Data of UD-CF-PA12 Laminates

Sample no.

Fmax (N)

C (ram~N)

a (mm)

W (mm)

B (ram)

Gnc (kJm - 2)

1 2

470 458

3'7.10 -3 4-I. 10 -3

12-1 15"2

9"4 9'9

2-7 2.7

2"7 3-5

conditions, the result was surprisingly high (G~c = 5.0 kJ m - 2). It is, however, in quite good agreement with the intralaminar fracture energy as calculated from the Kc test of the 1-90°]8 CT samples (G¢=5"9kJm-2). When comparing these data with interlaminar fracture energies given for other highly tough thermoplastic matrix systems, e.g. C F - P E E K composites 22 - 24 the value is found to be of the same order of magnitude, though somewhat higher. But also the PA-12 matrix itself has, for example, in terms of the notched Charpy fracture energy, aN, a much higher value (PA-12: a N = > 4 0 - 3 4 0 k J m -2, E , ~ I . 2 G P a ; x4'15 PEEK: a N = 5 4 k J m -2, E ~ 3"7 GPaI*). In fact, when looking at the fracture surfaces of broken DCB specimens, a remarkable degree of plastic matrix deformation is visible (Fig. 17), which is higher than has been observed for corresponding PEEK-matrix samples (Gic PEEK= 2"0 kJ m - 2).22 - 24Mode II interlaminar fracture energies as derived from two different E N F samples are listed in Table 2. The average value of Gzi~= 3.1 kJ m - 2 is also here higher than corresponding values achieved with a C F - P E E K system under similar external loading conditions (Gnc = l ' 6 k J m - 2 ) . 25"26 The reason is, again, based on the tougher matrix system used here. This is reflected in the same way on the mode II fracture surfaces (Fig. 18). Here, the matrix between the fibers has been highly deformed under shear loading. In both cases, i.e. mode I and mode II, the fibers between the deformed matrix tips are barely covered with matrix material, which is again a clear indication of the very poor bonding between the components of this composite system. For better comparison between the two deformation modes, micrographs of two other fracture surfaces are depicted in Fig. 19.

4 CONCLUSIONS This study was carried out to demonstrate the following: (a)

That continuous carbon fiber bundles, interspersed with fine thermoplastic PA-12 powder and surrounded by a sheath of the same polymer, can be consolidated into different laminate forms.

118

K. Friedrich, T. Gogeva, S. Fakirov

(b)

That the flexural properties are in agreement with what would be expected from the theory for different laminate structures. Measurements must, however, be carried out with recommended rather than arbitrary specimen geometries. (c) That through-thickness fracture toughness properties generated for different laminates reflect the tough nature of the matrix system used, and the effect of poor interfacial bonding which results in high K c values for composites with 0 ° fibers transverse to the crack front. (d) That interlaminar fracture energy data obtained for mode I and mode II crack-opening conditions are rather high when compared to those of other thermoplastic matrix systems. Their values are not so surprising, however, when the extensive deformation of the matrix material on the fracture surface is taken into account. As a final comment, we note that the data given in this paper may only be considered as preliminary data without any relevance for future users of this material. The reasons for this statement are: (1) a very limited a m o u n t of material was available at the time this study was carried out; (2) a rather primitive consolidation procedure was used for the preparation of the laminates; (3) there was a strict limitation on specimen geometry which did not allow us to stay within the recommended dimensions for the different test samples used. The data reported can, however, be seen as indicators of what might be expected with respect to the toughness profile of composite laminates made from this particular fiber/matrix combination.

ACKNOWLEDGEMENT Support of this project by a contract of research cooperation between the G e r m a n Department of Research and Technology and the Bulgarian Government is gratefully acknowledged (BMFT 227-9211-BUL).

REFERENCES 1. Brandt, J. & Richter, H., Hochleistungsverbundwerkstoffe mit thermoplastischer Matrix. Kunststo[Je, 77 (1987) 1,404. 2. Friedrich, K., Carlsson, L. A., Smiley, A. J., Walter, R. & Gillespie, J. W., Mechanisms for rate effects on interlaminar fracture toughness of carbon/epoxy and carbon/PEEK composites, J. Mater. Sci., July (1988) submitted for publication. 3. O'Brfidaigh, C. M. & Mallon, P. J., Effect of forming temperature on the properties of polymeric diaphragm formed APC-2 components. In Proc. 2nd

Thermoplastic impregnatedfiber bundles

4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16.

17. 18.

19. 20. 21. 22.

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Technical Conf. American Soc.for Composites, Newark, Delaware, USA, 23-25 September, 1987, pp. 36-46. Smiley, A. J. & Pipes, R. B., Simulation of the diaphragm forming of carbon fiber/thermoplastic composite laminates, Ibid. pp. 218-21. Eveno, E., Gillespie, J. W., Schultz, J. M. & Vinson, J. R., Welding of thermoplastic composites: numerical and experimental investigation. Ibid., poster presentation. Beyeler, E. P., Phillips, W. A. & Giiceri, S. I., Experimental investigation of laserassisted thermoplastic tape consolidation, J. of Thermoplastic Composite Materials, 1 (1988) 107. Cogswell, F. N., The Processing of Thermoplastic Structural Composites, Imperial Chemical Industries (ICI), Wilton, UK, 1987, internal report. Stolze, R., Aramid and carbon fiber reinforced thermoplastics. Proc. 25th Int. Man Made Fibres Congress, Dornbirn, Austria, 24-26 September, 1987. Lang, R. W., Stutz, H., Heym, M. & Nissen, D., Polymere HochleistungsFaserverbundwerkstoffe. In Proc. GDCh-Conference: Neue Polymere-Spezielle Eigenschaften-Moderne Technologien, Bad Nauheim, FRG, 14-15 April 1986. Atochem Report, Thermoplastic Polymers in Powders for Composites, ATOCHEM, Paris, 1985. Carlsson, L. A. & Pipes, R. B., Experimental Characterization of Advanced Composite Materials, Prentice Hall, Englewood Cliffs, N J, 1987. Carlsson, L. A., Gillespie, J. W. & Trethewey, B. R., Mode II lnterlaminar fracture of graphite /epoxy and graphite/PEEK. J. Reinforced Plastics and Composites, 5 (1986) 170. Hull, D., An Introduction to Composite Materials, Cambridge University Press, Cambridge, 1981. Elias, H. G. & Vohwinkel, F., Neue polymere Werkstoffe fiir die industrielle Anwendung. Hanser Verlag, Miinchen, FRG, 1983. Hellerich, W., Harsch, G. & Haenle, S., Werkstoff-Fiihrer Kunststoffe. Hanser Verlag, Miinchen, FRG, 1983. Bradley, W., Fracture toughness of polymer matrices in relation to interlaminar fracture toughness of composite materials. In Application of Fracture Mechanics to Composite Materials, ed. K. Friedrich. Elsevier Science Publ., Amsterdam, in preparation. Friedrich, K., Fractographic analysis of. polymer composites. Ibid. Lang, R. W., Heym, M., Tesch, H. & Stutz, H., Influence of constituent properties on interlaminar crack growth in composites. In High Tech--the Way into the Nineties, ed. K. Brunsch, H. D. G61den, C. M. Herkert. Elsevier Applied Science Publ., Amsterdam, 1986, pp. 261-72. Friedrich, K., Fracture mechanical behavior of short fiber reinforced thermoplastics. Fortschritt-Berichte VDI-Zeitschriften, VDI-Verlag, Dfisseldorf, FRG, Reihe 18, Nr 18, 1984. Slepetz, J. M. & Carlson, L., Fracture Mechanics of Composites, A S T M STP 593, Philadelphia, 1975, pp. 143 54. Friedrich, K., Hornbogen, E. & Sandt, A., Ultra-high strength materials. Fortschritt-Berichte VDI-Zeitschriften, VDI-Verlag, Dfisseldorf, FRG, Reihe 5, Nr 82, 1984. Gillespie, J. W., Carisson, L. A. & Smiley, A. J., Rate dependent mode I interlaminar crack growth mechanisms in graphite/epoxy and graphite/PEEK. Composite Sci. Technol., 28 (1986) 1-15.

120

K. Friedrich, T. Gogeva, S. Fakirov

23. Donaldson, S. L., Fracture toughness testing of graphite/epoxy and graphite/PEEK composites. Composites, 16 (1985) 103-12. 24. Crick, R. A., Leach, D. C., Meakin, P. J. & Moore, D. R., Interlaminar fracture morphology of carbon fibre/PEEK composites. J. Mater. Sei., 22 (1987) 2094-104. 25. Gillespie, J. W., Carlsson, L. A. & Pipes, R. B., Finite element analysis of the end notched flexure specimen for measuring Mode II fracture toughness. Composite Sci. Teehnol., 27 (1986) 177-97. 26. Chapman, T. J., Smiley, A. J. & Pipes, R. B., Rate and temperature effects on Mode II interlaminar fracture toughness of composite materials. In Proc. I C C M VI/ECCM 2, Elsevier Applied Science Publ., London, July 1987, Vol. 3, pp. 3.295-305.