JOURNAL OF POLYMER SCIENCE: Polymer Physics Edition VOL. 13. 1177-1186 (1975)

Tensile Properties of Ultradrawn Polyethylene

NUMA J. CAPIATI* and ROGER S. PORTER, Polymer Science and Engineering, and Materials Research Laboratory, University of

Massachuselts, Awtherst, Massachusetts 01 002

Synopsis High-density polyethylene filaments prepared by a solid-state deformation in an

Instron capillary rheometer show unusually high crystal orientation, chain extension, axial modulus, and ultimate tensile strength. The Young's modulus and ultimate tensile strength have been determined from stress-strain curves. Gripping of this high modulus polyethylene has been a problem heretofore, but the measurement of ultimate tensile strength has now been made feasible by a special gripping procedure. Tensile moduli show an increase with sample preparation temperature and pressure. Values as high as 6.7 X 10'' dyne/cm2 are obtained from samples extruded a t 134°C and 2400 atm and tested a t a strain rate of 3.3 X 10-4 sec-1. The effect of strain rate and frequency on modulus has also been evaluated by a combination of stressstrain data and dynamic tension plus sonic measurements over nine decades of time.

INTRODUCTION

The mechanical properties of semicrystalline polymers are sensitive to the and state of ~rientation.~" In fact, the high modulus and tensile strength of commercial synthetic fibers are due to morphological transformations and to chain orientation procedures. The unconventional drawing process used here to induce orientation and chain extension in- volves extrusion of crystalline polyethylene in an Instron capillary rheome- ter a t draw ratios to over 300: 1. The draw ratio is calculated as the ratio of rheometer reservoir to capillary cross-sectional area. This study thus deals with the tensile properties of ultradrawn polyethylene strands pro- duced by a deformation principally in the capillary entrance region of an Instron capillary rheometer.

EXPERIMENTAL Tensile experiments at constant rate of strain were carried out on ultra-

oriented polyethylene samples extruded a t different temperatures and pressures. Stress-strain curves up to the break point were obtained, for the first time, on such polyethylene strands.

The material used was exclusively high-density polyethylene, Du Pont Alathon 7050, M , = 58,000 and M , = 18,000. An Instron testing machine

* On leave from Universidad Nacional del Sur, Bahia Blanca, Argentina. 1177

0 1975 by John Wiley & Sons, Inc.

1178 CAPIATI AND PORTER

model TTM was used. The strain was measured by means of a strain gage extensometer, with a magnification of 2000 : 1 for moduli and 200 : 1 for tensile strength determinations.

The gripping of samples for tensile tests in the Instron machine has been a problem heretofore because of the anisotropy and low friction coefficient of these high-modulus filaments. The best results were obtained by hold- ing the sample ends in a sandwich-like assembly. From the outside, the grips are: two aluminum blocks with axial, semicircular grooves of the same diameter BS the strand, two pieces of silicon carbide sand paper (2p grain size), and, in the center, the strand. The whole system was fitted in the machine crosshead by special jaws. This arrangement prevented slippage out of the jaws and allowed these samples to be broken in the middle for the first time.

The tensile modulus was determined by the tangent to the stress-strain curve at a strain level of 0.03%.

Strands were prepared using a variation of the method proposed by Southern and P ~ r t e r . ~ Constant pressure and temperature were main- tained through the two stages of the process: ( 1 ) Crystallization from the melt at elevated pressure without extrusion. Flow is arrested by setting a conically tipped steel needle to plug the capillary exit. (2) Solid-state extrusion of the already crystallized material through a specially designed, capillary with a small entrance angle. A brass, stepdown capillary was used with 30 to 10" included entrance angle regions of 1.3 and 0.5 cm lengths, respectively, an 0.14cm diameter and a 1.0-cm length. The draw ratio was 46.

The strand quality and length are sensitive to the geometry and inner surface smoothness of the capillary. Lubrication with Teflon spray proved to be useful. The Teflon adheres to the capillary wall forming a layer of extremely low friction coefficient that improves the solid-state extrusion process. Successive filaments extruded without the further application of Teflon spray provided strands of similar length and properties. Capil- laries with low entrance angles give longer strands. Filaments about 20 cm long can be extruded, free of defects, with a final included entrance angle of 10".

Strands were prepared at extrusion pressures of 1270-2400 atm and temperatures from 130 to 136°C. Temperature and pressure were con- stant in a given extrusion to within zkO.5"C and f0.250j,, respectively.

Young's moduli are averages of from three to five repeat experiments. The confidence interval, for a reliability of 0.95, is f5Q/,.

RESULTS AND DISCUSSION

Table I shows typical values of elongation to break, tensile strength, and Young's modulus together with the corresponding literature values for carbon steel, aluminum, and glass fibers. It can be seen that the poly- ethylene filaments are comparable in rigidity with glass fibers and alumi-

ULTRADRAWN POLYETHYLENE 1179

TABLE I Comparison of Mechanical Properties

Tensile Tensile modulus, strength, Elongation dyne/cm2 d y ne/cm* to break,

Material x 10-11 x 10-9 % Carbon steel, SAE 1020” 20.8 4.5-6.3 25-32 Aluminum- 6.9 0.9-1.7 15-45

Ultra-oriented high-density Glass fiberb 6.9-13.8 3.9-6.9 4 .8-5.4

polyethylene 6 . 7 4.8 3 . 2

a J. H. Perry, Chemical Engineering Handbook, McGraw-Hill, New York, 4th ed..

b N. J. Parrat, Fiber Reinforced Materials Technology, Van Nostrand, New York, 1972, 1969, Section 23, p. 31.

p. 68.

Fig. 1. Effect of strain rate and solid-state extrusion conditions on tensile modulus.

num; the steel modulus is about three times higher than our ultra-oriented polyethylene. The strand tensile strength is similar to that of steel and glass fiber and considerably greater than that of aluminum. The prop- erties in Table I are stated on the usual volumetric basis; if they were expressed per unit mass, the figures for the polyethylene would be much higher, in particular when compared against steel.

The influence of preparation and test conditions on the mechanical characteristics of the filaments are summarized in Figure 1. The modulus increase with both extrusion pressure and temperature. lo The effect of strain rate is less pronounced. As can be seen, the change with strain rate

1180 CAPIATI AND PORTER

is small in spite of the considerable expansion of the modulus scale. behavior is as expected for a highly crystalline, highly oriented material.9

This

Effect of the Extrusion Pressure

The effect of preparation pressure on the mechanical properties of the fibers can be explained on considering both the pressure dependence of crystal size and the molecular mechanism of crystal deformation.

A change in the pressure of extrusion will affect mainly the initial crystal- lization step in the preparation process. Pressure is applied sufficiently fast that pressure equilibrium is achieved before significant crystallization takes place. The equilibrium melting temperature increases almost linearly with pressure.l1

AT = T , - T , + O.O2P,

where AT is the undercooling, T , the equilibrium melting point for a per- fect polyethylene crystal, T , the crystallization temperature, and P , is the crystallization pressure in atmospheres. Consequently, from the equation, undercooling increases in proportion to pressure. Since the lamellar crys- tal thickness is nearly independent of pressure up to 2500 atm" and the crystallization rate does not change considerably up to 5000 atm,12 the main effect of pressure on the original morphology is considered to be the in- crease in undercooling resulting from the shifting of the equilibrium melting point.

A basic characteristic of nucleation theories is the importance of under- cooling on the critical dimensions of the nuclei.13

I* = ~ u ~ T J A H ~ A T

p* = ~ T U , ~ T , ~ / A H ~ ~ A T ~

where I * is the critical dimension of the nucleus parallel to the chain (c axis), p* is the critical dimension normal to the c axis, uc and u,, are the interfacial free energies corresponding to faces perpendicular and parallel to the c axis, respectively, AHr is the heat of fusion per unit volume, T , is the absolute melting temperature, and AT the undercooling. The change in nucleus dimensions with undercooling is proportional to the change in the equilibrium crystal size. l4

At a crystallization temperature of 134°C the polyethylene undercooling for 1270 and 2400 atm are, respectively, 33 and 56°C." Assuming that the interfacial free energies and the heat of fusion remain constant in this pres- sure range,12 the crystal dimensions for the pressure limits are 1*1270 = 1.6 1*2400 and p*~270 = 2.6 p*2400. The stronger dependence of the lateral di- mension on undercooling promotes a significantly more elongated crystal at 2400 than a t 1270 atm.

The relation between thc crystal dimensions and the macroscopic me- chanical properties of the oriented strands can be analyzed considering the molecular model proposed by Peterlin.1s16e16 This model explains the de-

ULTRADRAWN POLYETHYLENE 1181

Fig. 2. Tensile fracture surface.

formation mechanism of a scmicrystallinc polymer under tensile load and is consistent &ith considrrablc morphological and mechanical information published on drawn synthetic fibers. It is, essentially, a combination of the folded chain and fringed micelle concepts.

When a semicrystalline polymer with a spherulitic structure is subject to an axial strain, a three-step deformation process results: (1) plastic d c formation of the original spherulitic structure; (2) transformation from spherulitic to fibrillar structure by lamellar opening, micronecking, and subsequent extension to form microfibers ; and (3) plastic deformation of the fibrillar structure.

The fractional chain extension within the microfibers makes them ex- tremely strong and rigid; the macroscopic properties are yet considerably lower. The lateral adhesion between fibrils is limited to weak van der Waals’ forces. In consequence, slippage between bundles of microfibers is likely the common failure mechanism rather than chain backbone break- age.17-19 The tensile fracture surface of a polyethylene strand, Figure 2, shows that the rupture results from shear rather than tensile deformation, as can be evidenced from the needlelike fracture surfaces.

The spherulite material, as crystallized in the first stage of the prcpara- tion piocess, is transformed into fibrils by extrusion through the capillary. The cross-sectional area of the resulting microfibers will be proportional to the width of the lamellar crystals from which thcy were formed. Rlor- phologies crystallized a t higher pressures were shown to contain smaller

CAPIATI AND PORTER

1 ' I I I , I

I200 1400 1600 1800 2000 2200 2400 E X T R U S I O N PRESSURE ( o t m )

Fig. 3. Effect of extrusion pressure on tensile modulus.

and more elongated crystals. Thus, a better chain packing can be expected at higher pressure and the resistance to slippage will also be increased since the fibril-to-fibril closeness provides more anchorage points among them. The mechanical' properties will be improved in consequence. Such fea- tures are consistent with experimental results. Figure 3 shows an increase of about 60% in modulus when the extrusion pressure is doubled.

Effect of Extrusion Temperature

The orientation process involving the transformation from spherulitic to fibrillar structure requires considerable deformation of both the crystalline and amorphous phases. The lamellar crystals are oriented with their c axis parallel to the extrusion direction, allowing, a t higher strains, the for- mation of the fibrillar structure. l.ls

Molecular mobility will increase with solid-state extrusion temperature. Therefore, higher temperatures below the melting point will increase the deformability thus improving the organization of the fibrillar structure. The probability for tension concentrations in the final strand will also be lower. This is consistent with an increase in modulus and tensile strength with extrusion temperature. The narrow temperature range covered is due to experimental limitations. At less than 130°C the polyethylene can start crystallizing even before the increased pressure is established, whereas a t over 135"C, relaxation and partial melting of the fibrils occurs at the capillary exit. Consistentwith this, Figure 4 shows a decrease in modulus above 134°C for these sample diameter and processing conditions. Since the maximum in the modulus-temperature curve is controlled by the rate

ULTRADRAWN POLYETHYLENE 1183

/ STRAIN RATE

PRESSURE: 2100 a m

I I I 1

130 131 132 I 33 134 135 EXTRUSION TEMPERATURE (‘C)

Fig. 4. Effect of extrusion temperature on tensile modulus.

of cooling of the strand, it may be expected to depend on those variables. It was found, for example, that a transparent filament can be obtained at temperatures up to about 150°C if air is blown at the capillary exit to increase the cooling rate of the fiber!

Correlation Between Transient and Dynamic Properties

Because of the remarkable magnitude of these moduli from stress-strain experiments of constant strain rate, it is desirable to develop comparisons with independent measures of modulus as by dynamic and sonic means. The comparison of data by different methods is also the usual procedure to study the modulus behavior over an extensive range of time when the desired range cannot be covered by a single method of measurement.

Dynamic moduli were obtained by Rheovibron (Toyo Measuring Instru- ments Inc.) measurements on the ultraoriented polyethylene samples. This method has been used previously on such samples. lo Precautions are taken to provide reliable values for such high-modulus polymers. The sonic modulus was measured in a Dynamic Modulus Tester PPM-5R (H. M. Morgan Co., Inc.). The authors wish to gratefully acknowledge the valuable collaboration of Mr. S. Colantoni for carrying out the sonic measurements. The single sample was an ultraoriented strand extruded at 135°C and 2100 atm (see Fig. 5).

The difference betm-een the dynamic storage modulus E’(w) and the Young’s modulus E(t) should be small a t t = l / w . The difference can be related through the relaxation spectrum as Marvinm pointed out. Further-

1184 CAPIATI AND PORTER

Fig. 5. Correlation among transient, dynamic, and sonic modulus. 1

more, Catsiff and TobolskyZ1 have proposed a simplification of the Marvin equation.

~ ~ ( w ) l w - l / t - E(O = H(7)*(m)(r-t (1)

where H ( 7 ) is the spectrum of relaxation times T ; *(m) = (s/2) csc (mr/2) - r(m); I' is the gamma function; m is the negative slope of a logarithmic plot H ( 7 ) versus 7.

For a highly crystalline polymer it can be assumed that 0 < m < 1 be- cause of the relative flatness of the relaxation spectrum curve.22 Further, it can be shown that23 *(m) 6 0.58.

The right side of eq. (1) can be shown to be only a few percent of the storage dynamic modulus value in the range of frequencies used.

The expression for H ( 7 ) as a function of E'(w) is:24

For 0 < m < 1 results sin(ms/2)/(ms/2) 6 1. The value of d[log E'(w) I/ d[log w ] is about 0.06 in the range of frequencies to be Then, the relaxation spectrum will be, at most, 6% of the dynamic modulus value. Hence, the right side of eq. (1) is less than 3.5% of E'(w) . This minor deviation thus falls within experimental error.

Figure 5 shows the modulus behavior through nine decades of frequeney (or strain rate). Young's storage dynamic and sonic modulus values are given. In general, the results are in agreement with stress-strain data taken at comparable strhins, about 0.03%.

In the strand preparation process, the fiber length depends upon the capillary geometry. It has thus not been easy to get a single sample which fits the geometrical requirements of both transient and dynamic modulus test methods. Different strand diameters and slightly different methods of preparationlo were used for the dynamic tests. Reasonable

ULTRADRAWN POLYETHYLENE 1185

agreement among the data on different samples and by different methods is apparent in Figure 5. The relatively weak linear dependence shown in Figure 5 is as expected for a highly crystalline and highly oriented polymer.

CONCLUSIONS

For the solid-state extrusion process for polyethylene : I . The tensile modulus was found to increase with the preparation pres-

sure. The dependence is consistent with the improved fibril-to-fibril ad- hesion and better chain packing which arise from higher preparation pressures.

Increased preparation temperatures below the melting point facili- tates the formation of the fibrillar structure by increasing the material de- formability during extrusion. Strands produced at higher temperatures have a higher modulus and tensile strength probably because of a decrease in localized tension concentrations.

The slope of a plot of E’(w) versus w is small for these oriented, highly crystalline polyethylenes. The Young’s modulus data thus can be related directly, within the experimental error, a t t = l/w. The agreement is good among the modulus measurement methods.

We report here the first stress-strain modulus data on ultradrawn polyethylene morphologies of such high strength. Interesting alternate approaches have also produced high-modulus, highly drawn, linear poly- ethylenes.26~27

2.

3.

4.

This study has been carried out by a fellowship from the Consejo Nacional de Inves- tigaciones Cientificas y Tecnicas de la Republica Argentina granted to N.J.C. and in part by a grant from the National Science Fmndation, GK 41268, and by support of the Office of Naval Research, contract N00014-67-A-0230-0011.

References 1. A. Peterlin, Textile Res. J., 42, 20 (1972). 2. I. Sakurada and K. Kaji, J. Polym. Sci., C31. 57 (1970). 3. T. T. Wang, J. Appl. Phys., 44, 2218 (1973). 4. J. Pate1 and P. J. Phillips, J. Polym. Sci., B11, 771 (1973). 5. F. C. Frank, Proc. Royal Soc. London, Ser. A , 319, 127 (1970). 6. D. W. Hadley, P. R. Pinnock, and I. M. Ward, J. Muter. Sci., 4,152 (1969). 7. S. Maruyama, K. Imada, and M. Takayanagi, Int. J. Polym. Mater., 1,211 (1972). 8. G. Meinel and A. Peterlin, J . Polym. Sci., A-2, 9, 67 (1971). 9. J. H. Southern and R. S. Porter, J. Macromol. Sci.-Phys., B4, 541 (1970).

10. N. E. Weeks and R. S. Porter, J. Polym. Sci., Polym. Phys. Ed., 12,635 (1974). 11. B. Wunderlich and T. Arakawa, J. Polym. Sci., Pt. A , 2,3697 (1964). 12. P. D. Calvert and D. R. Uhlmann, J. Polym. Sci., Pt. A d , 10, 1811 (1972). 13. L. Mandelkern, Crystallization of Polymers, McGraw-Hill, New York, 1964, p.

14. Ibid., pp. 296, 322. 15. A. Peterlin, J. Muter. Sci., 6,490 (1971). 16. A. Peterlin, in Advances in Polymer Science and Engineering, K. D. Pae, D. R.

17. M. Kurokawa and T. Ban, J. Appl. Polym. Sci., 8,971 (1964).

248.

Morrow, and Y. Chen, Eds., Plenum Publishing Corp., New York, 1972, pp. 1-19.

1186 CAPIATI AND PORTER

18. T. Hinton and J. G. Rider, J . Appl . Phys., 39,4932 (1968). 19. R. E. Robertson, J . Palm. Sci., Pt. A-8, 7,1315 (1969). 20. R. S. Marvin, Phys. Rev., 86, 644 (1952). 21. E. Catsiff and A. V. Tobolsky, J . Colloid Sci., 10,375 (1955). 22. J. D. Ferry, ViseOelastic Properties of Polymets, 2nd ed., Wiley, New York, 1970,

23. Ibid., p. 106. 24. M. L. Williams and J. D. Ferry, J . Polyn. Sei., 11,169 (1953). 25. N. E. Weeks, personal communication. 26. G. Capaccio and I. M. Ward, Polymer, 15.233 (1974). 27. A. G. Gibson, I. M. Ward, B. N. Cole, and B. Parsons, J . Ma&. Sci., 9, 1193

p. 64.

(1974).

Received September 26, 1974 Revised January 16, 1975