Effects of rubber content and matrix structure on static and fatigue fracture in ABS copolymers

363

EFFECTS OF RUBBER CONTENT AND MATRIX STRUCTURE ON STATIC AND FATIGUE FRACTURE 1N ABS COPOLYMERS M. RINK, F. BRIATICO-VANGOSA Politecnico di Milano, Milano, Italy L. CASTELLANI Enichem, Mantova, Italy

ABSTRACT Rubber toughening mechanisms in acrylonitrile-butadiene-styrene (ABS) may be affected by the dispersed phase size, structure and content and by the characteristics of the styrene-acrylonitrile copolymer (SAN) matrix, namely by its molecular weight distribution and acrylonitrile (AN) content. In this work a series of ABS samples having different matrix average molecular weights and AN contents, and different dispersed rubber-phase content were prepared. For all materials time to fracture initiation and crack speed as a function of the stress intensity factor were determined under cyclic loading, while J-resistance curves were determined under static loading. Directly compared, the results obtained indicate that fatigue fracture is positively affected by molecular weight, but scarcely or, sometimes, negatively by the rubber content; on the contrary, static J-resistance is higher for higher rubber contents, but is not significantly affected by matrix molecular weight. However, a proper consideration of the time scale in which fracture occurs in the two loading modes, and of the strain energy release rates involved in fatigue testing, shows that these results are only apparently contrasting.

KEYWORDS acrylonitrile-butadiene-styrene (ABS), fatigue, crack initiation, crack propagation, J-testing, the styreneacrylonitrile copolymer (SAN) molecular weight, acrylonitrile (AN) content.

INTRODUCTION Interest in fatigue testing of polymers arises not only from the growing use of plastics in structures subjected to cyclic loading, but also from the particular and sometimes new insights which can be gained, through this type of test, on the deformation and fracture processes. Fatigue loading has been reported to superimpose additional material responses and consequences on those commonly encountered during "static" loading or deformation [1 ]. Fatigue failures involve an initiation stage followed by a propagation stage [2-5], and specific material variables may have different effects on each of them [6]. Increasing molecular weight has been reported to have large positive effects on fatigue resistance for many polymers [1,7], even though values of static fracture toughness varied little over the same molecular weight range.

364

M. RINK, F. BRIA TICO- VANGOSA, L. CASTELLANI

In rubber-modified thermoplastics, second-phase content and structure have been found to bear different and more problematic effects on fatigue resistance than on fracture toughness as measured during monotonic loading [8,9]. Two main experimental approaches to the investigation of fatigue resistance of polymers can be found in the literature: the "stress-number of cycles to failure (S-N) curve" method, in which an unnotched specimen is subjected to controlled cyclic loading and the number of cycles to crack initiation, Ni, and to final specimen failure, Nf, are recorded; and the "fatigue crack propagation (FCP)" approach (or "Fracture Mechanics approach"), where cyclic load is imposed on a notched specimen and the crack length, a, is continuously monitored, this allowing the measurement of the crack propagation rate, da/dN, and its correlation to the instantaneous value of the stress intensity factor range at the crack tip, AK. Fatigue in rubber modified glassy polymers was studied through the "S-N-" approach by J.A. Sauer and C.C Chen [6,9,10], who found that while high impact polystyrene (HIPS) deforms by craze nucleation and growth, ABS under fatigue deforms primarily by shear deformation, with crazing developing only at a later stage. Particles in ABS were found to cavitate and/or to be bypassed by the advancing crack, whereas particles in HIPS are mostly fractured by the propagating fatigue crack. Both initiation and propagation were investigated: rubber content in HIPS and ABS was found to increase the resistance to crack propagation, but to decrease the resistance to crack initiation. C.B. Bucknall and T.A. Faitrouni [11,12] investigated fatigue crack propagation (FCP) in SAN and ABS with varying rubber contents and SAN molecular weights. They reported a significant reduction in FCP rates as a result of addition of rubber up to 7.5%, but little effects after further addition of dispersed phase. This was compared with the fracture toughness Gic, under monotonic loading, which was found to increase up to about 10 % rubber, and with the Charpy notched impact strength, steadily increasing up to 25% dispersed phase. Rubber effects on FCP were also found to be smaller for higher SAN molecular weights. Fatigue resistance was found to markedly increase with increasing molecular weight in pure SAN and, to a lesser extent, in ABS. This was interpreted, with reference to a model proposed by Michel and Hertzberg [7,14], in terms of the fraction of molecules longer than the critical molecular weight for entanglements, Me. Effects of the dispersed phase content on rubber toughened SAN were also studied in [ 13,4]. Fatigue crack propagation resistance in ethylene-propylene-diene rubber reinforced SAN [ 13] increases up to about 25% rubber after which it strongly decreases; resistance to static fracture initiation, JIc, and propagation, dJ/da, in ABS [4] were found instead to reach a plateau above 15% and 35% rubber content respectively. Aim of the present work is to further investigate the fatigue fracture of ABS, giving particular consideration to the different roles played by the SAN matrix and by the rubber-phase in fatigue as compared to non-cyclic fracture testing. We chose to focus the attention on two characteristics of the SAN matrix, molecular weight and composition (AN wt. %), and on the rubber content. The weight average molecular weight, Mw, of SAN was considered with reference to the entanglement molecular weight, Me, which is known to depend on the AN content [15], in order to compare samples having different Mw/Me ratio, instead of simply different Mw. This was made on the basis of the well documented (see for example ref. [16]) key role played by parameters Me and Mw/Me in the interpretation of deformation micromechanisms in glassy polymers. Mechanical response of the materials was evaluated by means of FCP tests and J-resistance curve measurements.

365

Effects of Rubber Content and Matrix Structure

MATERIALS ABS samples were prepared by melt mixing, in a twin screw extruder, SAN resins having various AN contents and average molecular weight with SAN-grafted Polybutadiene (PB) particles obtained by emulsion polymerisation. Diameters of PB particles are distributed between 0.1 and 0.4 ]am, and the grafted SAN to PB weight ratio is 0.62. Table 1 lists, for each sample, the PB weight content and the SAN matrix characteristics. These are the AN weight content, the weight average molecular weight, Mw, and the Mw/Me ratio (see above). As can be observed, the samples are representative of two well-defined and separate levels for each of the three structural parameters to be investigated: matrix Mw/Me ratio (levels 7.5 and 10.6), matrix AN wt. % (levels 24 and 33) and rubber-phase content (levels 10 and 25). Samples are labelled with three letters, the first indicates the AN wt. % (a=24% or A=33%), the second the rubber-phase wt. % (b = 10% or B=25%) and the third the Mw/M~ ratio (m=7 or M = 10). Specimens for mechanical testing were machined from 210xlT0x6 mm plates obtained by compression moulding at 185 ~ with a pressure of 5 MPa. Table 1. Materials Sample

Mw

Mw/M~

aMB

AN in SAN wt.% 24

123000

11.2

PB wt.% 25

amB

24

82000

7.4

25

AMB

33

94000

10.2

25

AmB

33

70000

7.6

25

AMb

33

94000

10.2

10

Amb

33

70000

7.6

10

EXPERIMENTAL Tensile tests. Tests were carried out at room temperature on dumb-bell shaped specimens (ASTM D638M-82, type M-II), by means of a screw-driven Instron dynamometer, at a crosshead displacement rate of 10 mm/min. A 12.5 mm clip gauge extensometer was used to measure the strain on the specimens. Fatigue tests. Testing was performed at room temperature, by means of a hydraulically driven Instron machine operating in load control mode, in accordance with the ESIS/TC4 draft protocol proposed in 1997 [17]. The loading wave was sinusoidal, of tension-tension type, at a frequency of 1 Hz and an R-ratio (minimum load/maximum load) of 0.3. Single edge notched (SE(T)) specimens (width = 35 mm, length = 100 ram, thickness = 6 mm) were used. Notches were introduced by alternatively sliding a sharp blade (tip radius about 10 lam). Initial notch length values, a0, were in the range 2.5-4 mm, the exact value for each specimen being measured after the test on the fractured surface. Crack propagation was monitored by means of a video camera, equipped with a magnifying lens (magnification 20 - 30 x), connected to a video tape recorder. Printed marks on the specimen surface regularly spaced along the expected crack path, allowed the measurement of the crack length after the test on the recorded images.

366

M. RINK, F. BRIATICO-VANGOSA, L. CASTELLANI

Non-cyclic fracture tests. J-resistance curves were determined at room temperature in accordance with the 1992 ESIS/TC4 protocol [ 18], in three-point bending, on SE(B) specimens having dimensions 120x25x12.5 mm. Notching was performed with a sliding sharp blade, as above described for the fatigue tests, up to a crack length to specimen width ratio, a/W, of 0.6. Crosshead speed was 1 mm/min.

RESULTS Tensile tests

Tensile modulus E was measured as the slope of the initial linear portion of the stress-strain curves and the yield stress, Cry,was taken as the maximum of the same curves. Average values and standard deviations obtained after testing five specimens for each of the samples are reported in Table 2. As can be expected, the rubber-phase content has a strong effect on tensile modulus and yield stress, with lower values corresponding to higher rubber-phase levels. The observed mechanical properties appear not to be significantly affected by the matrix structure (AN content and molecular weight) when the rubber content is high (25%), while at comparatively lower rubber contents (10%) rigidity and yield stress are greater for higher molecular weight of the matrix. Table 2. Tensile and static fracture tests results Sample Tensile modulus (MPa)

Yield stress (MPa)

J0,2

(dJ/da)Aa=0.2

(kJ/m 21

(kJ/m 31

aMB

1690 + 30

36.3 + 0.5

5.1

12.8

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1700 _+ 100

37.0 +_0.4

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12.7

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1730 + 50

39.2 _+0.5

5.5

14.1

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1660 +_40

36.9 _+0.3

6.2

13.9

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2900 + 100

57_+ 1

5.3

10.5

Amb

2500 _+80

53 _+ 1

5.0

12.4

Fatigue tests

Two stages were considered when examining the results of fatigue testing: crack initiation and crack propagation. Direct measurement of the number of cycles for crack initiation, Ni, from the video recording is difficult because of the very low crack speed at the beginning of propagation. Crack length, a, was therefore plotted against the number of cycles, N, after the test, and a conventional value for the number of cycles to crack initiation, N0.2, was taken at 0.2 mm crack advance, as interpolated on the experimental a vs. N curve. A check of the reliability of this method was made by performing on identical specimens, a series of tests which were stopped at progressively increasing N values, after which the specimens were cooled in liquid nitrogen and broken at high loading rate. The extent of fatigue crack propagation could be measured on the resulting fracture surfaces, thus obtaining an independent set of a(N) values: this confirmed the existence of a crack initiation phase and the good correlation between N0.2 and Ni. Propagation stage data were analysed by means of the conventional double logarithmic plots of crack propagation rate daMN vs. the stress intensity factor range, AK. The latter was computed by means of the following geometry calibration equation [ 17]:

367

Effects of Rubber Content and Matrix Structure

AP

AK =

5.q~-~

1)

B 4 w (20 - 13~ - 7~ ~)~ in which AP is the load range within a cycle, B the specimen thickness, a the crack length, W the specimen width and o~- a/W. The crack propagation rate, daMN, was determined from the a(N) curves. In the following, fatigue test results are presented by plotting the applied stress intensity factor range, AK, as a function of the number of cycles at crack initiation, Ni, and, as far as crack propagation is concerned, by means of da/dN vs. AK log-log plots. The effects of the three structural parameters under investigation, i. e., matrix molecular weight, matrix AN content and rubber-phase content are considered separately. Molecular weight o f the matrix. Figure 1 compares the crack initiation resistance of low and high Mw/Me

samples in three different cases: high rubber content and low AN in SAN (fig. la); high rubber content and high AN in SAN (fig. l b); low rubber content and high AN in SAN (fig. l c). In all cases the higher molecular weight is clearly associated with greater Ni values. The effect appears smaller for the lower AN content (figs. la and lb), while about the same effect is observed at low and high rubber contents (figs. lb and 1c). Molecular weight effects on crack initiation resistance appear in all cases to become less imporrtant in the high AK range. The same low and high Mw/Me samples are compared during the crack propagation stage in figure 2: at low and high AN (with 25% PB) in figs. 2a and 2b; at low and high rubber content (with 33% AN in the matrix) in figs. 2b and 2c. An increased resistance to propagation (smaller crack speed da/dN) is generally observed for the high molecular weight samples in the low AK range. Similarly to what found for crack initiation resistance, when the stress intensity increases, the molecular weight effect is gradually reduced. The molecular weight effect is larger in the high AN-high PB sample (fig. 2c) than in the others. A N content in the S A N matrix. Samples with lowand high AN content are compared for crack initiation in

figure 3 for low and high molecular weights. All samples in fig 3 have high rubber content (25%). The AN content has a positive effect on Ni, particularly when Mw/Me is high and AK is low. A slight positive effect of AN content on propagation resistance is also found, as figure 4 shows for low (fig. 4a) and high (fig. 4b) molecular weights. 1,5t

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M. RINK, F. BRIATICO-VANGOSA, L. CASTELLANI

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Rubber-phase content. Figure 5 illustrates how the rubber content influences crack initiation for high A N content samples having low and high Mw/M~. A small positive effect results at low molecular weights, but no significant effect can be detected when the high molecular weight samples are compared. Rubber-phase effects on propagation are also not simple. Figure 6a shows that for low molecular weights the low rubber content sample is more resistant at high stress intensity, but seems to become less resistant in the low stress region. When Mw/Me is high (Fig. 6b) the low rubber content sample appears more resistant than the high rubber content one in the whole AK range under investigation. 1,5~ . . . . . . . . . . . .

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0,6

Effects of Rubber Content and Matrix Structure

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Fig. 5. Applied stress intensity factor range vs. number of cycles to fracture initiation: effect of rubber content. B corresponds to high rubber content.

Non cyclic fracture tests Also for the J-resistance curves resulting from the static fracture, the effects of the three investigated structural parameters are considered separately. Molecular weight o f the matrix. In figure 7 low and high M , J ~ samples are compared in the cases of high rubber content and low A N (Fig. 7a), high rubber content and high A N (Fig. 7b); low rubber content and high AN (Fig. 7c). Differences due to the molecular weight appear to be within the order of magnitude of the experimental uncertainty, with the possible exception of the high rubber-high A N case (Fig. 7b) where a slightly higher toughness is exhibited by the low Mw/M~ sample. 12/,,,!,,,,,,,, f l 9 amB

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1

370

M. RINK, F. BRIATICO-VANGOSA, L. CASTELLANI

ANcontent in the SANmatrix. Figure 8 shows the effect of AN content, on samples with a constant PB

content of 25% and low (Fig. 8a) or high (Fig. 8b) Mw/Me values. Increasing the AN content in SAN has a positive effect which is more pronounced when molecular weight is comparatively low. 1 2 / , , , ~,,, ~ , , , ! , , ,

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Aa (mm) Fig. 8. J-resistance curves: effect of SAN acrylonitrile content. A corresponds to high acrylonitrile content. Rubber-phase content. Fracture toughness is higher for the materials of higher rubber-phase contents, as Fig. 9 shows in the cases of low (Fig. 9a) and high (Fig. 9b) molecular weight, for samples having an AN content in SAN of 33%. In the low molecular weight case (Fig. 9a), J-Aa curves for the two rubber levels are parallel to each other over the experimental Aa range; for high Mw/Me the high rubber curve has a greater slope, with values equivalent to those of the low rubber sample at small Aa and considerably greater as the crack length increases. These features are conveniently represented by means of the pseudoinitiation value, J0.2, taken from the J-resistance curves at Aa = 0.2 mm and by the slope of the curve at the same point (dJ/da)Aa: 0.2 [18]. Values are reported in Table 2.

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371

Effects of Rubber Content and Matrix Structure

DISCUSSION Besides a general positive effect of the AN content in the SAN matrix, material response is quite different in the fatigue and static test conditions: - Mw/Me, which has a strong positive effect on fatigue initiation and propagation resistance (Figs. 1,2), exhibits, if anything, rather negative effects on static J-resistance 0~ig. 7). rubber content, which has a strong positive effect on static toughness (Fig. 9), has little or no influence on fatigue crack initiation (Fig. 5) and small, mostly negative effects on fatigue crack propagation. These two features will be separately discussed in the following. M~/M~. The dimensions of the specimens used to obtain the J-resistance curves (Figs. 7, 8 and 9) largely comply with the size requirements for size independence, i.e. specimen thickness and initial ligament length are greater than 25 J/~y [18]. Plane strain, small scale yielding conditions can therefore be assumed and an estimate of the stress intensity factor at crack tip, KI, reached during J testing can be obtained from the equation [ 19]:

K~

J.E (1_ v2)

2)

where E is the tensile modulus and v is the Poisson ratio assumed to be 0.35. Stress intensity factors applied to the materials in the two test conditions can therefore be compared. Static stress intensity factors at fracture initiation, K0.2, calculated from J0.2 by means of eq. 2, and converted into an "equivalent" AK by using the constant R=0.3 ratio used in fatigue testing, are plotted vs. time to fracture initiation in fig. 10, together with fatigue initiation data from Figs 1,3 and 5 (number of cycles, N, in fatigue is equivalent to time in seconds because of the constant frequency of 1 Hz used in all the tests). Crack initiation during J testing in all cases occurred at higher stress intensities and at shorter times than in fatigue testing.

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1000

1500

2000

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time to fracture initiation (s) Fig. 10. Applied stress intensity factor range vs. time to fracture initiation: static and fatigue data.

-4,5

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Fig 11. Crack speed vs. stress intensity factor range during crack propagation static and fatigue data.

372

M. RINK, F. BRIATICO-VANGOSA, L. CASTELLANI

Crack speeds in the static J tests can also be calculated by differentiating a plot of the measured crack extensions Aa against the relevant test time. This allows a comparison between static propagation and fatigue propagation data. Fig. 11 shows this in the case of samples Amb and AMb: K0.2 and Kmax in J testing (calculated from J0.2 and Jmaxby eq. 2) and again converted into AK are plotted vs. the relevant crack speeds in mm/s together with the fatigue crack propagation data from Fig. 2c (mm/cycle equivalent to mm/s). For clarity only best-fit fatigue curves are reported. As may be observed, the stress intensity factor-crack propagation rate relationship in the static tests is very similar to that found in fatigue testing. This has been proven to be the case also for the other samples. This indicates that once the crack starts to propagate, the stress intensity factor crack relation is determined mainly by the stress level rather than by the nature of the loading. The different material responses in the two test conditions can therefore be essentially ascribed to the different crack initiation conditions, lower stress intensities and longer times being easily obtained in fatigue testing, while the J-testing procedure makes the fracture start at shorter times and higher K. Plastic deformations processes, and particularly craze growth, are known to be time (and temperature) dependent [ 16]. Fracture in SAN occurs through craze growth and breakdown, and two mechanisms have been shown to be involved in craze growth: chain scission, at short times/low temperatures, and chain disentanglement, at long times/high temperatures [16]. Craze initiation stress is not molecular weight dependent when occurring through chain scission, but is strongly molecular weight dependent when the disentanglement mechanism is active [6]. The observed molecular weight effects, strong and positive in fatigue testing at low AK, but negligibly small in J testing and in fatigue at high AK, can be related to different craze initiation and growth times in the two tests. In fact, since crack initiation in J testing occurred at high K and short times (Fig. 10), craze initiation and growth times were also short and therefore in this case chain scission was likely to occur. In fatigue testing a much wider range of K was explored and therefore, crack and craze initiation and propagation could take place either at high K and short times where chain scission occurs (no molecular weight effect), or at low K and long times where disentanglement dominates (strong molecular weight effect). Rubber content. A direct comparison between J-resistance curves and fatigue log(da/dN) vs. log(AK) plots is not correct. J-resistance curves show the energy required for crack advancement, while the energy

information is not contained in the fatigue crack propagation plots, where the stress intensity-crack speed relationship is described. Energy can however be easily introduced in the representation of fatigue data, by using the strain energy release rate range AG in place of AK. AG was calculated in a first approximation by using the linear elastic K/G equivalence and a constant Young's modulus, equal to the value reported in Table 2. Fig. 12 shows the results obtained in the case of two samples with different rubber contents, AMB and AMb. The conventional log(da/dN) vs. log(AK) plot for these samples (Fig. 6b) shows a negative effect of the rubber content in the whole investigated AK range; when the same data are plotted vs. AG, an opposite result is obtained, with a strong positive effect of the rubber content on fatigue propagation resistance. The differences in "toughening efficiency" of the rubber phase in static and fatigue testing appear therefore to be an outcome of the data representation, and not the results of really different material responses in the two test conditions.

Effects of Rubber Content and Matrix Structure

373

'---' -0,5

-1,5

Z -2,5 "a e~ -3,5 o

tent

M

-4,5

-3

-2,6

-2,2

-1,8

Log [AG (kJ/m2)] Fig. 12. Crack speed vs. energy release rate range: effect of rubber content.

CONCLUSIONS Fatigue crack propagation tests and J-resistance curve determinations have been carried out on a series of ABS samples having various AN contents in SAN, SAN molecular weights and rubber-phase contents. In agreement with results previously reported in the literature, different structural parameters appear to be effective in the two (static and dynamic) loading conditions. However, after a careful analysis of the results we conclude that a direct comparison between conventional J-R curves and log(daJdN) vs. log(AK) plots, as far as the investigation of the effects of the materials' structural parameters on the fracture resistance is concerned, may be misleading. In fact: Fracture initiation in fatigue testing usually occurs at lower AK and longer times than in J testing. This can promote ductile deformation mechanisms in the matrix, which are more molecular weight dependent than the mechanisms dominating at shorter times and higher stress intensifies. Toughening efficiency of the rubber phase appears lower in fatigue than in static tests when the conventional log(da/dN) vs. log(AK) fatigue data representation is used. A different perspective is however obtained when log(da/dN) is plotted vs. log(AG). -

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REFERENCES 1. 2. 3. 4. 5. 6. 7.

Hertzberg, R.W., Manson, J.A. (1980). Fatigue of Engineering Plastics. Academic Press, N. Y. Williams, J.G. (1977). J.Mater.Sci. 12, 2525. Williams, J.G (1984). Fracture Mechanics of Polymers. Ellis Horwood Ltd.,Chichester, U. K. Rink, M., Imbrighi, D., Castellani, L. and Pavan, A. (1990). ECF 8 Fracture behaviour and design of materials and structure, D. Firnao Ed., EMAS Ltd., London,, p.201 Rink, M., Guidetti, B., Frassine, R., Castellani, L. (1994). J. Mater. Sci. 29,3071. Sauer, J.A., Chen, C.C. Adv. Polym. Sci. 91/92.(1990). H.H. Kausch Ed., Springer-Verlag, p.69. Michel, J.C., Manson J.A., Hertzberg, R.W. Polymer, (1984). 25, 1657.

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Manson, J.A., Hertzberg, R.W., Carlingo M.J., Hahn, M.T., Hwang, J., Turkanis J., Attalla, G. (1985). Int. Conf. "Toughening of Plastics II", 2-4 July, London, The Plastics and Rubber Inst. Sauer, J.A., Chen, C.C.- Polymer Eng. Sci. 1984, vol.24, n. 10, pp.786-97. Sauer, J.A., Chen, C.C. Adv. Polym. Sci. 52/53.(1983)., H.H. Kausch Ed., Springer-Verlag, p.169. baitrouni, T. A. (1990). Effects of structure on the fatigue behaviour of ABSpolymers, Ph.D. Thesis, Cranfield Inst. Of Technology, Bucknall, C.B. and Faitrouni, T. A. (1991). 8~ Int. Conf. "Deformation, Yield and Fracture of Polymers", Cambridge, 8-11 April. The Plastics and Rubber Institute. Ricc6, T., Rink, M., Caporusso, S., Pavan, A. (1985). Int. Conf. "Toughening of Plastics II'" 2-4 July. London, The Plastics and Rubber Inst. Michel, J.C., Manson J.A., Hertzberg, R.W. (1985). Sixth Int. Conf.. "Deformation, yield and fracture of polymers", Cambridge, U.K 1-4 April. The Plastics and Rubber Institute. Lomellini, P., Rossi, A.G. (1990) Makromol. Chem. 191,1729. Kramer, E.J., Berger, L.L. Adv. Polym. Sci. 91/92.(1990). H.H. Kausch Ed., Springer-Verlag, p. 1. European Structural Integrity Society (ESIS), Technical Committee 4 "Polymers and Composites" ."Test Method for Tension-Tension Fatigue Crack Propagation", 1997. European Structural Integrity Society (ESIS), Technical Committee 4 "Polymers and Composites"; "A Testing Protocol for conducting J-Crack Growth Resistance Curve tests on Plastics", May 1995. Anderson, T.L. (1995). Fracture Mechanics, Fundamentals and Applications. CRC Press.