Creep and creep damage of glass fibre reinforced polypropylene

Composite Structures 24 ( 1993) 233- 244

Creep and creep damage of glass fibre reinforced polypropylene J. Hugo, M. Sova & J. Ci~insk~ National Research Institute for Materials, Opletalova 25, 11312 Prague I, Czech Republic

Composites based on isotactic polypropylene matrix and short glass fibres were tested for creep, creep damage and creep rupture at different temperatures and stress levels. Four composites of 5-30 wt% of glass fibre content having standard fibre/matrix adhesion were used, while two materials with 20 and 30 wt% of fibres were prepared as elevated adhesion composites. Unreinforced polypropylene was tested as a reference material. The rate of steady-state creep (which was determined by an arbitrary method) was used as a criterion for the examination of glass fibre content and fibre/matrix adhesion effects on creep. Secondary criteria, type and extent of damage and time to rupture were applied. Up to 400C the steady-state creep and creep damage of all tested polypropylene composites are retarded by glass fibres according to the fibre content. At 60"C and above, there is a distinct difference between composites of standard and elevated fibre/matrix adhesion. At higher temperatures the creep and creep damage are more or less accelerated by higher content of standard adhesion fibres; these fibres induce crazing (as a typical creep damage) at relatively low stress. Fibres of elevated adhesion shift the processes of non-recoverable creep and crazing to the higher stress range, even if their craze inducing activity remains. It was found that the creep, creep damage and creep rupture are controlled by the steady-state creep mechanism of polypropylene, a conclusion which was deduced from the stress and temperature dependence of the creep rate, as well as from the morphological analysis of damaged and ruptured specimens.

1 INTRODUCTION T h e general reason for reinforcing thermoplastics by high modulus fibres lies in minimising the elastic and creep compliance, as well as in elevating the creep strength, while the long-time service of the material at temperatures near the matrix d e p e n d e n t upper limit is of special interest. T h e effect of randomly distributed glass fibres varies with the fibre/matrix volume ratio, with the fibre length distribution, with the quality of the interface, etc. A question rises: which processes of the deformation and damage are influenced by the fibres and to what extent? T h e practical aspect of this question is evident: manufacture and processing of such composites are costly operations and high efficiency can be reached only in the case where the composite brings (compared with the base polymer) mechanical i m p r o v e m e n t in the broad range of stress, temperature and time. In contrast to metallic ~,z and ceramic 3,4 materials, creep and creep rupture of plastics and polymer composites have not been (in the last 20

years since K a m b o u r and Robertson's contribution 5) comprehensively reviewed. T h e lack of critical considerations is evident, namely concerning the mechanisms of the non-recoverable part of the creep strain, as well as of the damage processes. T h e creep and creep fracture is mostly discussed among many other aspects of the mechanical properties of polymers. 6-s Isotactic polypropylene was tested for creep many times, starting with practice-oriented work of Turner in the 1960s. 9-12 In recent years the attention turned to polypropylene composites. 13-~5 As concerns the creep mechanism of polypropylene, the contribution of Ward and co-workers should be mentioned. 16 Their experiments fall under the systematic research of stress relaxation and creep of oriented polymeric structures. ~7 Deformation mechanisms responsible for the yielding of polypropylene are discussed by Porzucek and coworkers. 18 Recently, Di Liello et al. 19 published the results of yield stress testing of polypropylene/ glass fibre composites, comparing uncoated and aluminium coated fibres to the fibre content and

233 Composite Structures 0263-8223/93/S06.00 © 1993 Elsevier Science Publishers Ltd, England. Printed in Great Britain

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Hugo, M. Sova, J. ~i~insl~

interface effects. They try to correlate the fibre content and interface quality with the activation energy and activation volume in the sense of Eyring's theory. Simplifying the problem, the processing of experimental creep data can be divided into two areas. In the first, an analytical expression for the creep curve is proposed and proved. Surprisingly, some empirical equations fit the creep curves of polymers for extremely long time periods. 2° Within the second area, typical for the research of metallic materials, an approach based on the theory of thermally activated flow of Eyring is widely used. Creep rate and its temperature and stress dependence are the main means for the interpretation of results. 21-23 Exact determination of the true activation energy brings some higher experimental demands, the activation volume in polymeric structures has not yet been fully clarified.17 However, omnipresent troubles come from the viscoelastic response of polymeric structures. Total creep strain is a superposition of elastic, delayed elastic and non-recoverable strains. To reach the steady-state creep where the SherbyDorn curve, log(de/dt) versus e, is in its minimum in the broad range of stress and temperature, is more than an experimental problem. 21 For the distinction between recoverable and non-recoverable strain, creep and recovery are measured. 24 In the case of long-time creep experiments, where (besides theoretical) practical results are required, such an approach is questionable. Some help could be found in the fact that in the short-time period after loading the strain-time relation is exponential, while in the long-time period the strain increases with time according to the power law.22 In the latter case the power law relation between creep rate g and stress 0 is characteristic. Exponent values should reflect the mechanism of the creep. However, it has been well known for a long time that the steady-state creep of polymers is generally not a very stable state; some pre-rupture phenomena (crazing, etc.) are to be expected before tertiary creep and fracture appear.

process was used for the preparation of the series containing 5, 10, 20 and 30 wt% of glass (see Table 1). Materials having standard interface quality are designated SA. Two materials with elevated interface adhesion (designated EA) were prepared by the addition of UCARSIL PC into the compounded melt. After preliminary checking of specimens taken from compression moulded plates, injection moulded objects and/or small plates, injection moulding of standard shape specimens was finally chosen as a preferential method. The reason for this was the small variation of results of different testing methods within a large number of specimens. Injection moulded specimens were made under carefully controlled conditions on the CS 88/63 machine. The overall length of dumbbell samples according to ~SN 64 0605 was 150 mm, the cross-section was 10 mm x 4 mm; the dimensions of the beams (separately prepared for the dynamic testing) were 140 mm x 12 mm x 3 mm. All composites and samples were produced in and supplied by the technological laboratory of TIU Neratovice. The quality of the samples was checked by different morphological methods. A simplified description of the samples structure is: a thin surface layer has a random orientation of fibres, melt-flow direction oriented fibres are in the interlayer, random oriented fibres are in the centre. A typical well known fibre-length distribution was found: the mean length was approximately 0.4 mm. There was no observable difference between SA and EA specimens. 2.2 Tensile creep and modulus testing

A lever type, constant tensile load creep apparatus (40 units) was used for testing. Room tempera-

Table 1. Materials under test

Material Nominal

2 EXPERIMENTAL 2.1 Materials and specimens

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Creep and creep rupture of GFRP ture measurements were run in the laboratory equipped with the standard temperature and humidity control system (23°C, 65% r.h.), while elevated temperature experiments were performed on the machines having a controlled oven for maintaining the sample temperature within + 0.5°C. The multisection heating system of the oven (the power of the sections being independently adjustable) together with the electronic control kept the temperature of the specimen (inclusive of clamps) within the narrow limits. Induction type extensometers outside the oven were used for the measurement of the distance between clamps: their sensitivity was 0.01 mm. Increase of the 'between clamps' distance was transformed to the elongation of the working portion of the specimen using experimentally determined correction factors. Analogue strain data were digitalised and computer processed. Complex modulus data were acquired from the resonance frequency (35-80 Hz) of the vibrating cantilever beam using Bruel + Kjaer apparatus.

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For checking of the creep damage, thin (10/~m) slices were taken from damaged and/or broken specimens, cut in the long axis direction, parallel with the thickness direction. The Jung microtome with metal knife was used at laboratory temperature. The slices were observed and photographed in the Jenapol polarised light microscope. Details of crazes were observed using a SEM Jeol JSM 35 microscope.

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3 RESULTS AND DISCUSSION

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The real part of the complex modulus E' as a function of the fibre content and temperature (Figs 1 and 2) gives evidence of the slight (nearly unmeasurable) effect of the interface in the small stress/strain range, even at elevated temperatures. However, some signals come from the values of damping; EA composites exhibit lower damping at 60 and 80°C. 3.2 The rate of creep

In contrast to the primary creep (the rate of creep is defined as g=de/dt), the rate of which is

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decreasing with the time of loading, and to the tertiary creep (where ~ is increasing with time while cumulating damage until the rupture appears), the steady-state (or secondary) creep is a process of constant rate increase of the nonrecoverable strain. The rate of this process can be called 'minimum creep rate' ~c/m~n"

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An arbitrary procedure has been adopted for the determination of the minimum creep rate gc/min" From the experiment a curve in the form of total strain (e) versus log time (t) is stored. The recoverable part of the strain is defined by the extrapolated straight line:

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The e and t data from the short-time starting period of the experiment were taken for the determination of n and b (see Fig. 3(a)). Creep strain e c is computed subtracting e, from e within the whole period of the experiment and an arbitrary non-recoverable creep curve, e c versus t, as a set of points is stored. This curve is cleared from scattering using numerical methods and the minimum creep rate Ec/min is determined (Fig. 3(b)). The duration of the steady-state creep varied from 10 ° to several thousands hours, according to the conditions of the experiment. Due to the limited term of the programme the minimum creep rate has not been observed in some cases of low stress and temperature; then the minimal observed value was used. Where the tertiary creep appeared, ~c/min values were fully confident. 3.3 Stress and temperature dependence of the creep rate

Creep experiments were running according to the scheme: -- temperature r (°C): 23, 40, 60, 80 -- stress o (MPa): 3, 6, 9, 12, 15, 18 -- time t (h): to the rupture or to - 104 Stress levels were chosen with regard to the testing temperature and materials formulation. When possible, four stress levels for one formulation and temperature were applied; three stress levels was the necessary minimum. The upper limit of accurate total strain e measurement was approximately 6%. As expected, all points in the double logarithmic plot, log gc/m~nversus log o, lie on straight lines (see Figs 4(a)-(d)). Using the equation (also known as Norton law): gc/min = Born (2) an examination of the materials formulation influence on B and m can be done. Certainly, temperature dependence must be taken into consideration. Being nearly independent on the glass content and/or the interface quality, the slope m is ris-

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very well. Formally, EA fibres are responsible for shifting the 'frozen matrix behaviour' into the range of higher temperatures. With the exception of low fibre content composite (SA2), standard adhesion fibres are losing their reinforcing effect at 60°C. To make the results more lucid, a new variable is introduced: instead of the nominal stress o a normalised stress o/E is used, E being Young's modulus from dynamic measurements. The E values are related to the creep testing temperature. As shown in Figs 7 and 8, eqn (2) is valid and materials are sorted into groups according to the temperature and formulations. In the low temperature range (Fig. 7) the slope of straight lines is different according to the adhesion while in the high temperature range (Fig. 8) the creep rate is rising with the rising glass content of SA composites. A compact group of points appears for unreinforced polypropylene, low glass content SA composite and EA composites. Plotting comparable minimum creep rate values (respective to

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the chosen value of normalised stress) versus temperature and glass fibre content, a general picture of the materials formulation effect comes out (Figs 9 and 10). As concerns the creep, there is no retarding effect of glass fibres in SA composites; moreover, relating the creep to the elastic response, their effect could be called 'weakening'.

L o w glass content S A composite seems to be the only exception. 3.4 Creep damage and r u p t u r e As the creep experiments were running up to the stage of tertiary creep or to the rupture, it was

239

Creep and creep rupture of GFRP

possible to examine the relation between the critical time tb and the rate of the steady-state creep (Fig. 11). Time tb is either the time to rupture or the time of evident starting of the tertiary creep. Considering the scattering of points in Fig. 11 as a natural phenomenon for this type of testing, the relation is clear enough to suppose that mechanisms of the secondary creep control the damage. The general meaning of the crucial importance of the matrix in this type of composites is thus well supported.

A selection of damaged specimens was inspected by the optical microscopy as to the type and extent of damage. Figures 12-20 illustrate some examples chosen to support the following statements:

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(b) Fig. 12.

REM evidence of the craze type damage shown on light micrographs.

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Hugo, M. Sova, J. ~i~insk~

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Solitary crazes with the tendency to open as cracks (after 2400 h). Unreinforced polypropylene; 80°C; 3 MPa; low creep rate.

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High number of small crazes at the high rate of creep. Unreinforced polypropylene; 80°C; 9 MPa (after 0"9 h).

-- the prevailing damage is the craze perpendicular to the stress direction (Fig. 12); - - a f t e r the slow and long-time creep the crazes are solitary, well developed, with the tendency to open as cracks (Figs 13 and 18); -- a high number of small crazes is typical when the creep rate is high (Figs 14, 16, 17); 'whitening' of specimens could eventually

appear while small crazes are precursors of voids (mushroom structure)(Fig. 20); crazes start preferentially at the fibre/matrix interface even if the fibres are parallel with the stress direction (Figs 15 and 16); fibres diminish the non-uniform crazing dependent on the structural anisotropy of the unreinforced moulded body (compare Fig. 14 with Fig. 18);

Creep and creep rupture of GFRP

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Fig. 16.

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Crazes start at fibres (after 5200 h). SA2; 80°C; 3 MPa; low creep rate; tertiary creep.

High number of small crazes starting at fibres (after 6 h), high rate of creep. SA2; 80°C; 9 MPa.

generally, the craze inducing activity of fibres is independent of the interface quality (Figs 18 and 19).

4 CONCLUSIONS Polypropylene composites containing short glass fibres with the typical length distribution, low

mean aspect ratio and more or less random orientation is a material the mechanical behaviour of which is governed by the matrix. Concerning the non-recoverable creep, the following conclusions (( 1 )-(5)) can be drawn from the results described above. (1) There is a matrix-dependent difference in the matrix/fibre interaction at low (approximately up to 40°C) and high (approximately above 500C) temperature. This may be due to two creep

J. Hugo, M. Sova, J. (7ginsk~

242

Fig. 17.

Joining of crazes, large extent of local plastic strains (after 26 h). SA8; 80°C; 6 MPa; high creep rate.

Fig. 18.

Solitary crazes and cracks (after 5700 h). EA8; 80°C; 9 MPa; low creep rate.

mechanisms of polypropylene; ~6 one or both are active according to the temperature and to the stress. Naturally, the freedom for the non-recoverable flow is very low in the glassy state. That is why in the low temperature range (even above Tg of the amorphous phase) the reinforcing effect of fibres in SA composites is not much different from better bound fibres, and why the low tem-

perature creep is retarded depending on the fibre content. (2) The mechanism(s) of the non-recoverable creep of polypropylene remains unchanged in the presence of glass fibres even if the fibres are well bound. The differences between the creep of polypropylene and that of polypropylene composites are only quantitative. The creep rate of the com-

Creep and creep rupture of GFRP

Fig. 19.

Evidence of the craze inducing activity of glass fibres (after 26 h). EA8; 80°C; 15 MPa.

Fig. 20.

Mushroom structure -- whitening before rupture (after 310 h). SA8; 80°C; 4.5 MPa.

posite depends evidently on the actual stresses in the matrix. Where the matrix/fibre load transfer is efficient, the creep of the composite is low. This effect is manifested in the high temperature range where the well bound fibres reduce the matrix flow in their vicinity and remain active as loadbeating elements. (3) Due to their length distribution and random orientation, the fibres can act as reinforc-

243

ing elements as well as potential defects causing additional local stresses and strains. If the matrix flow is predominant (high stress, high temperature), insufficiently bound fibres (even if they are long enough and well oriented) lose their reinforcing role and the majority of fibres turn into potential defects, accelerating the creep and damage. In such a case the concentration of fibres (or interfibre distance) seems to be important. Concerning

244

J. Hugo, M. Sova, J. ~i~inskf:

SA composites, low fibre content materials exhibit lower creep and damage at elevated temperature than high fibre content composites. (4) All fibres (or interfaces) are potential initiators of crazes. Besides the dilatational nature of crazes, their origin should be considered as connected with local shear stresses, the mechanism being tightly joined with the mechanism of the non-recoverable creep. Where the creep rate is low, the growth of sporadic crazes is a preferential mechanism; in principle, this rule is not affected by the fibres. The high frequency of crazes corresponds with the high creep rate. The complex (strain, time and temperature dependent) effect of fibres may be the reason why the tensile yield testing (even if the temperature and testing speed dependent data are available) does not reflect well the actual formulations of short fibre composites, not mentioning mechanisms important for the practical service of such materials. (5) The formerly expressed opinion that the rate of the non-recoverable creep and time to rupture are in a simple relation is supported by experiments described above. The same mechanism controlling the creep and creep rupture is to be considered in the case of polypropylene composites. This work has been carried out in the course of optimising the quality and assortment of polypropylene composites. The authors believe that some knowledge obtained here could be applied to other thermoplastic matrix composites of different formulations. The work is a part of the project supported by the Chemopetrol, Prague and Czechoslovak Academy of Sciences.

REFERENCES 1. Cadek, J., Creep kovov~ch materifil6 (The creep of metallic materials). Academia, Praha, 1984 (in Czech). 2. Gittus, J. H., Creep, Viscoelasticity and Creep Fracture of Solids. John Wiley, New York, 1975. 3. Cannon, W. R. & Langdon, T. G., J. Mater. Sci., 18 (1983) 1. 4. Cannon, W. R. & Langdon, T. G., J. Mater. Sci., 23 (1988) 1. 5. Kambour, R. P. & Robertson, R. E., Mechanical properties of plastics. In Polymer Science, Vol. 1, ed. A. D. Jenkins. North-Holland, Amsterdam, 1972. 6. Turner, S., Mechanical Testing of Plastk~s'. Godwin, Harlow, Longman, New York, 1983. 7. Ward, I. M., Mechanical Properties of Solid Polymers. John Wiley, Chichester, UK, 1983. 8. Kinloch, A. J. & Young, R. J., Fracture Behaviour of Polymers. Applied Science Publishers, London, 1983. 9. Turner, S., Polym. Eng. Sci., 6 (1966) 306. 10. Sinclair, J. E. & Edgemond, J. W., J. AppL Polym. Sci., 13 (1969)999. 11. Dixon-Stubbs, P. J., J. Mater. Sci., 16 ( 1981 ) 389. 12. Kitagava, M. & Matsutani, T., J. Mater. Sci., 23 (1988) 4085. 13. Gupta, V. B. & Lahiri, J., J. Composite Mater., 14 (1980) 286. 14. Trantina, G. G., Polym. Eng. Sci., 26 (1986) 776. 15. H6ninger, H. & Reichelt, E., Plaste u. Kautschuk, 35 (1988) 322. 16. Duxbury, J., Ward, I. M. & Parsons, B., J. Mater. Sci., 22 (1987) 1215. 17. Sweeney, J. & Ward, I. M., J. Mater. Sci., 25 (1990) 697. 18. Porzucek, K., Lefebvre, J. M., Coulon, G. & Escaig, B., J. Mater. Sci., 24 (1989) 3154. 19. Di Liello, V., Martuscelli, E., Ragosta, G. & Zihliv, A., J. Mater. Sci., 26 (1991) 2100. 20. Findley, W. N., Polym. Eng. Sci., 27 (1987) 582. 21. Sherby, O. D. & Dorn, J. E., J. Mech. Phys. Solids, 6 (1958) 145. 22. Kubfit, J., Rigdahl, M. & Seld6n, R., J. Appl. Polym. &:i., 20 (1976)2799. 23. Wilding, M. A. & Ward, I. M., Polymer, 22 ( 1981 ) 870. 24. Mindel, M. J. & Brown, N., J. Mater. Sci., 8 (1973) 863.