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On the dissolution behaviour of extended chain polyethylene fibres
European PolymerJournal VoL 17. pp. 157 to 161
0014-3057/81/0201-0157502.00/0
© Pergamon Press Lid 1981. Printed in Great Britain
ON THE DISSOLUTION BEHAVIOUR OF EXTENDED CHAIN POLYETHYLENE FIBRES J. C. M. TORFS, G. O. R. ALBERDA VAN EKENSTEIN and A. J. PENNINGS Department of Polymer Chemistry, State University of Groningen, Nijenborgh 16, 9747 AG Groningen, The Netherlands
(Received 6 May 1980) Abstract--The influence of stress on the dissolution behaviour of extended-chain high molecular weight
polyethylene fibres in p-xylene was investigated. Freely suspended in the solvent, the fibres dissolved at 119.5°, a temperature close to the equilibrium solubility temperature of 118.6° for perfect polyethylene crystals. However, when a stress of 0.4 GPa was exerted by straining the fibre 0.7~, it could withstand a temperature as high as 130° for at least three days. At still higher temperatures the induced stress relaxed completely, and dissolution immediately followed. These phenomena indicate that the fibre has a network structure. The cross-links are of a physical nature. Molecules are connected by topological defects such as entanglements, intertwinings and twist disclinations. These defects are trapped in crystallites; therefore the theory of Gee and Flory is applicable predicting that in such a system dissolution temperature of extended chain crystallites increases with stress. The required stress is transduced by tie molecules bridging the amorphous regions between crystallites. A study of dissolution under stress seems to be a direct method for the detection of topological defects such as entanglements.
tremely low rate of dissolution due to a specific structure. This implies furthermore that the influence of stress on the dissolution behaviour of polymers might constitute the basis for a new m e t h o d of investigation of fibre structure. Some preliminary results concerning polyethylene surface-growth fibres are reported in this paper.
INTROD UCTION In previous papers [1-3] a technique for continuous growth of fibrillar crystals of high strength and modulus has been described. Crystal growth in the polymer chain direction was accomplished by subjecting a fibrous seed crystal, in a solution of high molecular weight polyethylene (Mw = 4 x 106), to a flow field. The flow field was established by having a cylindrical stirrer rotate in a couette apparatus. Essential for the formation of the strongest fibres is that the seed crystal slides very close to the surface of the stirrer. This technique, referred to as the surface growth process, showed the remarkable p h e n o m e n o n of fibrous growth at temperatures as high as 123 °, which is well above the equilibrium dissolution temperature of 118.6 ° I-4, 5] of perfect polyethylene crystals. This p h e n o m e n o n is apparently closely related to a n increase in crystallization a n d dissolution temperature, to a temperature higher than the equilibrium dissolution temperature, as a result of stress in the growing fibre. Stress is related to the entropy decrease caused by the deformation. Earlier experiments have indeed shown that the fibre dissolved at these higher temperatures when the stress, applied by means of a take-up device, was released. This behaviour is similar to that of polyvinylalcohol fibres which, as mentioned by Frenkel [6-8], are less soluble when stress is being applied. This a p p a r e n t thermal stability is however an unexpected p h e n o m e n o n , because the application of stress to a fibre will cause an increase in its enthalpy [9] a n d free energy, so that the crystals in the fibre must become thermodynamically less stable, a n d should dissolve at a lower temperature. Therefore the loss of solubility of the stressed fibres, at a temperature higher than their equilibrium dissolution temperature, must be accounted for by an ex-
EXPERIMENTAL
Linear polyethylene, Hi-fax 1900 from Hercules Inc. with weight average molecular weight of 4 x 106, was used in all experiments. The dissolution experiments with very little stress or no stress at all were carried out with a 0.2 m long fibre. This fibre was produced by the surface-growth technique [1, 2] at 115.5°; it had a cross-section of 1.39 x 10 - 9 m 2 as determined by its weight (while assuming a density of 103 kg/m3), and a tensile strength of 3.2 GPa, determined at a strain rate of 4 8 ~ per min. The fibre was submerged in the solvent, p-xylene, in a glass tube with an i.d. of 4 mm. A glass tube of such narrowness was used to prevent the fibre from curling. The sample was heated from 20 to 100° at a rate of 0.5 ° per min, and from then onwards at a rate of 0.25 ° per min. The change in length of the fibre during heating was measured with a cathetometer. In cases where little stress was desired, a metal wire weighing about 20 mg was connected to the fibre. The dissolution experiments involving considerable stress were carried out with fibres grown at 113.9° having a crosssection of 0.48 x 1 0 - 9 m 2 and a tensile strength of 2.7GPa. Four fibres, 0.3 m in length, were connected parallel by knotting them together at the ends. They were placed in p-xylene in a thermostated stress-relaxation apparatus described earlier [10]. The bundle of four parallel fibres was next rapidly strained until a stress of 0.4GPa was reached, corresponding to a strain of about 0.7%. From then on, the length of the fibre was kept constant. After the stress had during 3 hr been allowed to relax, the solvent was heated at a rate of 1° per min up to the preselected temperature. This temperature was maintained within _+0.2°.
157
J.C.M. TORFS,G. O. R. ALBERDAVAN EKENSTEINand A. J. PENNINGS
158
shrinkage(%) 15 10
-5 -10
'
100
110
=T(°C) 120
Fig. 1. Percentage of shrinkage vs temperature for a fibre heated in p-xylene in a 4 mm wide glass tube at a rate of 0.5°/min up to 100° and then at 0.25°/min. O: Upper curve: no stress. O: Lower curve: a stress of 0.4 MPa. The percentage of shrinkage is defined as A//lo x 100, where Al is the decrease of the length with respect to lo, the length at room temperature. Differential scanning calorimetry was carried out by means of a Perkin-Elmer DSC-2. The thermograms were obtained from a polyethylene fibre submerged as 4% (w/w) in p-xylene with the antioxidant dibutyl-p-cresol in sealed stainless steel pans [1 l]. The fibre was either kept at constant length by winding it on an iron wire, or it was free to shrink, because it was cut into 3 mm long pieces. Corrections for thermal lag were carried out according to the DSC manual, and indium was used for temperature calibration. RESULTS First of all, we examined the possibility of applying the equilibrium dissolution temperature of polyethylene in p-xylene of 118.6 ° [4, 5] to the fibres. In these experiments a fibre, to which no stress was applied and which was kept in a glass tube, was dissolved. When heated, such a fibre became transparent at 90°; at higher temperature, it shrank by slightly curling up to 16% at 119.4 ° as shown in Fig. 1 (solid circles). At 119.5 °, the fibre broke at several places and dissolved completely within a few minutes. These experiments proved that the unloaded fibre dissolves very dose to the equilibrium dissolution temperature of polyethylene in p-xylene of 118.6 °. When a fibre was submerged rapidly in p-xylene at 124 °, it dissolved visually within one minute, indicating the high rate of dissolution. The dissolution of unloaded fibres was also investigated by DSC of pieces of fibre, short enough to be able to retract easily. Therefore, pieces of fibre, 3 m m long, were heated in p-xylene. The ends of the endotherms were considered to be the dissolution temperatures. At heating-rates of 0.31, 0.62 and 5.0 ° m i n - 1 , the dissolution temperatures were 121.5, 121.7 and 124.0 ° respectively. The lowest value is comparable to the value of 119.4 ° obtained above at a rate of 0.25 ° per min. Simultaneous with dissolution, shrinking occurs (Fig. 1) due to the adoption of the molecules of a
more coiled conformation, first of all in the most defect regions. This seems to suggest that prevention of this coiling by applying a large enough stress to the fibre would reduce its solubility. When a small stress of 0.4 MPa was applied to a fibre, it dissolved in an essentially similar way. It started shrinking at about 100 ° (Fig. 1, open circles). At 119.0 ° it became transparent and elongated, and at 119.5 ° it broke and dissolved within a few minutes. It is remarkable that so small a stress (1/8000 of the tensile strength of 3 . 2 G P a at room temperature) is able to elongate the fibre. Elongation began as soon as the fibre, probably due to swelling, began to grow transparent. This indicates that the mobility of the long polymer molecules is very high in the swollen fibre, probably because crystallinity has disappeared. Dissolution experiments involving considerable stress were carried out with a bundle of four fibres. They were strained to obtain the initial stress and then kept at constant length. After 3 hr of stress relaxation in order to achieve equilibrium, the bundle was heated up to the preselected temperature, which was kept constant. The stress on the fibre from the moment the temperature was constant, is plotted in Fig. 2 for experiments carried out at various temperatures. (Stress was plotted instead of relaxation modulus because preliminary experiments indicated it to be the more fundamental parameter. This point will be further investigated in future.) At temperatures up to 132 °, the experiments revealed that the stress decreased most rapidly during heating and the first hours after temperature equilibrium had been established. Stress relaxation probably occurs because of taut tie molecules---which bridge disordered regions---slipping out of the crystaUites. The same process takes place in the case of the annealing of cold drawn fibrous material [12]. A few hours after reaching temperature equilibrium, the stress relaxed more slowly, and seemed to approach a limiting value. Thus the fibre seemed not to dissolve, and could still be seen after two days, even at 132.0 ° i.e. 13° above the equilibrium dissolution temperature. Experiments at 130.8 ° were carried out during 6 days and even then the fibre did not dissolve completely. Partial dissolution was investigated by weighing the fibres before and after the experiments. It was found in an experiment carried out at 130.8 ° that only (7 + 4)% of the fibre had dissolved after 3 days. This loss of weight might be the result of the existence of a separate fraction (e.g. a skin) dissolving more easily than the main part of the fibre; the possibility of very gradual dissolution however, cannot be rejected as a possible explanation at present. These results indicate that, up to 132 °, the fibre does not dissolve or does so only extremely slowly. At the higher temperature of 136.5 °, the stress relaxed fully within 7 hr. Then complete dissolution followed within a few minutes, indicating again that the stability of the fibre results from the stress. At temperatures between 132 and 136 °, stress relaxation takes place more slowly and is irregular because of differences in stress within the bundle of fibres, causing different dissolution times. Stress relaxation, followed by dissolution, still takes hours even at 136.5 °. Therefore a quickly heated
Dissolution of extended chain PE fibres 50
159
STRESS (MPa)
z.O
1
3(3
.
•
i
116 5°CL 1250 °C
1308 °C 132 O°C
lO
o
0
500
1000
1500
1 2000
TIME(ram) I 2500
Fig. 2. Stress plotted vs time for a polyethylene fibre in p-xylene at the indicated temperatures. The equilibrium dissolution temperature is 118.6C. The fibre was previously heated at a rate of l:,/min. I Pa = 10dyn/cm 2 = 1.45 x 10 4 p s i = 1.02 x 10-Skgf/cm 2 = 1.13 x 10-ag/denier(ifd= I).
sample should dissolve at a considerably higher temperature; this effect could be investigated by DSC. In these experiments, contraction of the fibre was prevented by winding it on an iron wire. Heating at 5 ° per rain resulted in an endotherm with a maximum at 150°. This indicated that the fibre was partially crystalline up to above 150 °, i.e. far above the equilibrium dissolution temperature of 118.6°!
DISCUSSION
The dissolution experiments most remarkably show the rapid dissolution of the free fibre even at 119.5 ° whereas, when considerable stress is being exerted, the fibre can withstand 132 ° for at least 40 hr. At 132 ° the temperature is higher than the equilibrium dissolution temperature; therefore, the free energy of dissolved random coils is lower than that of molecules in the fibre (this effect is even larger, because the free energy of the fibre has increased by the strain energy). Thus the fibre is metastable and its apparent thermal stability arises from a very low or zero dissolution rate, for kinetic reasons. This low rate could be ascribed to the stress, but one must also consider other possibilities. Slow dissolution might be due to the well-known decrease of rate of dissolution with increasing crystallinity [13, 14] and with increasing molecular weight, caused by the formation of entanglements (these cause an enormous increase in the thickness of the surface gel-layer through which the diffusion of solvent has to take place [-13, 15, 16]). The fibres investigated in this study are highly crystalline (about 8._0%)and consist of high molecular weight polymer (My, = 4 x 106). But there are indications that the dissolution of the fibres is not always so slow. 1. The experiments in which no stress was exerted, and which were carried out with essentially the same materials as the experiments with stress, revealed a
visual dissolution of the fibre within a few minutes at 119.5 °, and within 1 min at 124 :. 2. Differential scanning calorimetry of unconstrained 3 mm pieces of fibre submerged in p-xylene revealed that, at the scan speed of 5 ° per mira the final melting temperature is only 2.5 ° higher than at the lowest rate (0.31 ° per rain). Hence a dissolution time of only 30 sec is estimated. Therefore, the extended chain fibres do not have an intrinsically low rate of dissolution: they dissolve slowly due to the stress. In previous work [17] it was shown that the fibres consist of parallel microfibrils, themselves sequences of alternating crystallites and amorphous regions [12, 17, 18]. The thermal stability of the stressed fibres prompts two further conclusions concerning the arrangement of their chain molecules. (i) The surface of the microfibrils does not act as a less stable region, where dissolution can start due to peeling the molecules off one by one. This requires the chains on the surface to be firmly anchored inside the crystalline blocks. In other words: "The chains traverse the fibril in lateral direction on account of topological chain defects such as intertwinings, twist disclinations [19], jogs [20], loops and entanglements". (ii) The microfibrils do not break-up in the amorphous regions, in which the segmental mobility must be high due to the liquid state and to swelling. "Therefore subsequent crystalline regions have to be connected by tie molecules" bridging the amorphous zone. This conclusion is in agreement with the concept of "continuous crystal" introduced by Porter [21], because the tie molecules between crystalline regions may crystallize near room temperature. These structural features can be depicted qualitatively in a schematic figure (Fig. 3 [22]), which shows a stressed fibre at about 130 °. The existence of topological defects is not surprising in view of the mechanism of formation of these fibres in the surface growth technique [1,2]. This technique is based on the stretching of an entangled
160
J. C. M. TORFS, G. O. R. ALBERDA VAN EKENSTEIN and A. J. PENNINGS
Fig. 3. Schematic representation of molecules in a stressed fibril submerged in a solvent at a few degrees above equilibrium dissolution temperature. Extended chain crystals are partially interrupted by disordered regions [bridged by tie molecules (a)], where the concentration of defects is higher and where cilia (b) emanate the fibril. The defects depicted are trapped entanglements (c), intertwinings (d), jogs 1-20] (e), and chain ends (f). network, followed by rapid crystallization. Hence, entangiements and intertwinings occurring in the solution are trapped in the fibril. The question now is: why do not the amorphous regions grow at the expense of the crystals? This would transform the fibril into a wholly amorphous state in which entanglements can slip off, and the structure disintegrate. The answer is that the dissolved tie chains (which are in equilibrium with the crystallites) have very low entropy due to the stress restricting their number of possible conformations [23, 24]. Thus transition of a crystalline segment to the stressed amorphous state is unfavourable. Hence, dissolution of the metastable fibres does not take place, because the fibres would have to dissolve via a path involving a very high free energy barrier: a large part of a crystallite would have to swell for a moment, in such a way that the entangled chains could slip through and release, and dissolution could start at the crystallite surface. Something else adding to the free energy barrier is the relatively high polymer concentration in the dissolved parts of the fibril. Thereby, the lowering of the melting point of the pure polymer by the solvent [4,25] from 145.5° to about 118.6° [4, 5] in very dilute solutions, is not complete in the vicinity of the melting crystal face. This behaviour is distinct from the dissolution of lamellae. Then the local polymer accumulation will spread by the unfolding of the chains and by the diffusion of entirely dissolved molecules.
At this stage we can conclude that the fibre structure strongly resembles a chemically cross-linked network. Due to the topological defects, the structure amounts to an entanglement network. The entanglements cannot vanish because the chains cannot slip out of the crystallites. Therefore the theory as formulated by Gee and Flory [26-28] for chemically crosslinked networks is applicable. From this theory it can be deduced that, by application of a macroscopic stress, the melting temperature of extended chain .crystallites rises. The stress is transduced to all crystallites by the tie molecules. The behaviour of the strained fibres is different above 132 °. Stress relaxes completely and then dissolution occurs as in the case of free fibres. Stress-relaxation indicates that the mobility in the crystals is so high that tie molecules can slip away. It should be noted that this limit of 132 ° closely approaches the limit for fibrous growth by means of the surface growth technique, which uses the same solvent. When the cross-sections of the fibres are plotted vs crystallization temperature and an extrapolation towards zero cross-section is made, the achieved temperature is about 130° [29]. This suggests that growth would be impossible at 130°. The temperature-limit thus obtained for the surface growth technique might very well arise from too rapid slippage of chains past each other, resulting in rapid creep. The stress relaxation temperature is also remarkably close to 130 °, where the lamellar thickness of melt-crystallized or long time annealed solutiongrown polyethylene increases by a factor of 2-5 [30]. This rather sudden increase in lamellar thickness suggests a higher chain mobility, dependent on temperature but not on a solvent (since the experiment was carried out without solvent). "The consistency of the three temperatures (stressdissolution: 132°; surface growth: 130°; annealing: 130°: suggests a sudden higher mobility for polyethylene chains in the crystalline phase." This is probably due to thermally generated defects, already significant at 110 ° [31]. Finally it should be remarked that the dissolution temperature of 119.5° of free fibres must be considered high. As a result of the high free energy of the crystal surface and defects, a dissolution temperature below the equilibrium value of 118.6 ° would be expected. Dissolution of polyethylene has indeed been observed within the range of 110-111 ° for high molecular weight lamellar crystals [4] and for high pressure crystallized extended chain crystals [32-34]. However, the behaviour of stacked extended chfiin crystallites connected by tie molecules (Fig. 3) must be different e.g. from re-entering chains in lamellae. In the former case dissolution takes place at a higher temperature because the entropy of the dissolving tie chains is relatively low, mainly because of the limitations on the number of conformations. Limitations result from the volume occupied by the crystallites and from the fixation of the tie chains in the crystallites [23, 24] they connect. Work is now in progress on other materials and on such aspects as the influence of initial strain and solvent quality on the dissolution process of these extended chain fibres.
Dissolution of extended chain PE fibres CONCLUSIONS The fibres produced by surface growth [1-3], containing extended chain crystallites, dissolve at a higher temperature than e.g. lamellar crystals. This effect is most pronounced when the fibres are subjected to a stress. Then the apparent dissolution temperature rises to 1 3 above equilibrium dissolution temperature (in p-xylene). The thermal stability is indicative of: 1. Tie molecules bridging the a m o r p h o u s regions between crystallites in a microfibril. 2. Topological defects such as entanglements. Thus the structure a m o u n t s to a network of entanglements which are trapped by crystallites. Crystallites do not dissolve, because the entropy of the stressed tie chains in the dissolved state is much lower than the entropy of free molecules in solution. Above 132 the fibres dissolve, because the applied stress relaxes, due to high chain mobility in the polyethylene crystals.
REFERENCES
1. A. Zwijnenburg and A. J. Pennings, Colloid Polym. Sci. 254, 868 (1976). 2. A.J. Pennings and J. C. Toffs, Colloid Polvm. Sci. 257, 547 (1979). 3. A. Zwijnenburg and A. J. Pennings, J. Polym. Sci, (Left.) 14. 339 (1976). 4. J. F. Jackson, L. Mandelkern and O. C. Long, Macromolecules I, 218 (1968). 5. A. J. Pennings, J. M. M. A. van der Mark and A. M. Kiel. Kolloid Z.u.Z. Polym. 237, 336 (1970). 6. S. Ya. Frenkel, S. A. Agranova. V. G. Baranov. Yu. N. Panov, T. I. Samsonova, V. S. Aldoshin and L. N. Korjavin. Preprints of the IUPAC Symposium on Macromolecular Chemistry. p. 159, Prague (1965). 7. S. Ya. Frenkel, V. G. Baranov, N. G. Belnikevitch and Yu. N. Panov. Vysomolek. Soedin. 6, 1917 (1964). 8, Yu. N. Panov, T. I. Samsonova. N. G. Belnikevitch, S. A. Agranova and S. Ya. FrenkeL Abstracts of papers presented at the 12th Annual Meeting of the Institute of Macromolecular Compounds, p. 49, Leningrad, 1965, 9. J. W. S. Hearle and E. A. Vaughn, In Polymer Science
10. l l. 12. 13. 14. 15. 16. 17. 18. 19. 20. 21. 22. 23.
24. 25. 26. 27. 28. 29. 30. 3 l. 32. 33. 34.
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(Edited by A. D. Jenkins) p. 877. Elsevier, North-Holland (1972). A. Posthuma de Boer and A. J. Pennings, Macromolecules 10, 981 (1977). C. J. M. Schouteten, S. Bakker, B. Klazema and A. J. Pennings, Analyt. Chem. 49, 522 (1977). A. Peterlin, Polym. Engn9. Sci. 18, 6, 488 (1978). D. W. van Krevelen, Properties of Polymers, pp. 439~.44. Elsevier, Amsterdam (1976). V. Kalisek and F. Krejci, Sb. Ved. Pr., Vys. Sk. Chemickotechnol., Pardubice 38, 105 (1978). K. Ueberreiter. In Diffusion in Polymer (Edited by J. Crank and G. S. Park). Academic Press, New York (1968). F. Asmussen, K. Ueberreiter and H. Naumann, In Diplomarheit. Fr. Univ. Berlin (1965). A. Zwijnenburg, P. F. van Hutten, A. J. Pennings and H. D. Chanzy, Colloid Polym. Sci. 256, 8 (1978). D. T, Grubb and A. Keller, Colloid Polym. Sci. 256, 218 (1978). J. C. M. Li and J. J. Gilman, J. appl. Phys. 41, 4248 (1970). B. Wunderlich, Maeromolecular Physics, Vol. 1, p. 450. Academic Press, New York (1973). R. S. Porter, A.C.S. Polym. Preprints 12, 2 (1971). A. J. Pennings, Makromolek. Chem. Suppl. 2, 99 (1979). H. G. Zachmann, Kolloid Z.u.Z. Polym. 216-217, 180 (1967). H. G. Zachmann, Kolloid Z.u.Z. Polym. 231, 504 (1969). P. J. Flory. J. chem. Phys. 17, 233 (1949). L. Mandelkern, Crystallization of Polymers, pp. 169 171. McGraw-Hill, New York (1964). G. Gee, Q. Rev. (London) 1,265 (1947). P. J. Florv, J. Am. chem. Soc. 78, 5222 (1956). A. J. Pennings, K, E. Meihuizen, In Ultra-high Modulus Polymers (Edited by I. Ward and A. Ciferri). Applied Science Publ. (1979). B. Wunderlich, Macromolecular Ph)sics, Vol. 1, p. 195. Academic Press, New York (1973). D. H. Reneker, J. Mazur and J. P. Colson, Preprints of Short Comm., 26th Int. Syrup. on Macromol., IUPAC, Mainz, Federal Republic of Germany (1979). A. Nakajima, F. Hamada, S, Hayoshi and T. Sumita, Paper presented at IUPAC Meeting, Tokyo-Kyoto, Japan (1966). B. Wunderlich and C. M. Cormier, J. phys. Chem. 70, 1844 (1966). C. M. Cormier and B. Wunderlich, J. Polym. Sci., A-2 4, 666 119661.
0014-3057/81/0201-0157502.00/0
© Pergamon Press Lid 1981. Printed in Great Britain
ON THE DISSOLUTION BEHAVIOUR OF EXTENDED CHAIN POLYETHYLENE FIBRES J. C. M. TORFS, G. O. R. ALBERDA VAN EKENSTEIN and A. J. PENNINGS Department of Polymer Chemistry, State University of Groningen, Nijenborgh 16, 9747 AG Groningen, The Netherlands
(Received 6 May 1980) Abstract--The influence of stress on the dissolution behaviour of extended-chain high molecular weight
polyethylene fibres in p-xylene was investigated. Freely suspended in the solvent, the fibres dissolved at 119.5°, a temperature close to the equilibrium solubility temperature of 118.6° for perfect polyethylene crystals. However, when a stress of 0.4 GPa was exerted by straining the fibre 0.7~, it could withstand a temperature as high as 130° for at least three days. At still higher temperatures the induced stress relaxed completely, and dissolution immediately followed. These phenomena indicate that the fibre has a network structure. The cross-links are of a physical nature. Molecules are connected by topological defects such as entanglements, intertwinings and twist disclinations. These defects are trapped in crystallites; therefore the theory of Gee and Flory is applicable predicting that in such a system dissolution temperature of extended chain crystallites increases with stress. The required stress is transduced by tie molecules bridging the amorphous regions between crystallites. A study of dissolution under stress seems to be a direct method for the detection of topological defects such as entanglements.
tremely low rate of dissolution due to a specific structure. This implies furthermore that the influence of stress on the dissolution behaviour of polymers might constitute the basis for a new m e t h o d of investigation of fibre structure. Some preliminary results concerning polyethylene surface-growth fibres are reported in this paper.
INTROD UCTION In previous papers [1-3] a technique for continuous growth of fibrillar crystals of high strength and modulus has been described. Crystal growth in the polymer chain direction was accomplished by subjecting a fibrous seed crystal, in a solution of high molecular weight polyethylene (Mw = 4 x 106), to a flow field. The flow field was established by having a cylindrical stirrer rotate in a couette apparatus. Essential for the formation of the strongest fibres is that the seed crystal slides very close to the surface of the stirrer. This technique, referred to as the surface growth process, showed the remarkable p h e n o m e n o n of fibrous growth at temperatures as high as 123 °, which is well above the equilibrium dissolution temperature of 118.6 ° I-4, 5] of perfect polyethylene crystals. This p h e n o m e n o n is apparently closely related to a n increase in crystallization a n d dissolution temperature, to a temperature higher than the equilibrium dissolution temperature, as a result of stress in the growing fibre. Stress is related to the entropy decrease caused by the deformation. Earlier experiments have indeed shown that the fibre dissolved at these higher temperatures when the stress, applied by means of a take-up device, was released. This behaviour is similar to that of polyvinylalcohol fibres which, as mentioned by Frenkel [6-8], are less soluble when stress is being applied. This a p p a r e n t thermal stability is however an unexpected p h e n o m e n o n , because the application of stress to a fibre will cause an increase in its enthalpy [9] a n d free energy, so that the crystals in the fibre must become thermodynamically less stable, a n d should dissolve at a lower temperature. Therefore the loss of solubility of the stressed fibres, at a temperature higher than their equilibrium dissolution temperature, must be accounted for by an ex-
EXPERIMENTAL
Linear polyethylene, Hi-fax 1900 from Hercules Inc. with weight average molecular weight of 4 x 106, was used in all experiments. The dissolution experiments with very little stress or no stress at all were carried out with a 0.2 m long fibre. This fibre was produced by the surface-growth technique [1, 2] at 115.5°; it had a cross-section of 1.39 x 10 - 9 m 2 as determined by its weight (while assuming a density of 103 kg/m3), and a tensile strength of 3.2 GPa, determined at a strain rate of 4 8 ~ per min. The fibre was submerged in the solvent, p-xylene, in a glass tube with an i.d. of 4 mm. A glass tube of such narrowness was used to prevent the fibre from curling. The sample was heated from 20 to 100° at a rate of 0.5 ° per min, and from then onwards at a rate of 0.25 ° per min. The change in length of the fibre during heating was measured with a cathetometer. In cases where little stress was desired, a metal wire weighing about 20 mg was connected to the fibre. The dissolution experiments involving considerable stress were carried out with fibres grown at 113.9° having a crosssection of 0.48 x 1 0 - 9 m 2 and a tensile strength of 2.7GPa. Four fibres, 0.3 m in length, were connected parallel by knotting them together at the ends. They were placed in p-xylene in a thermostated stress-relaxation apparatus described earlier [10]. The bundle of four parallel fibres was next rapidly strained until a stress of 0.4GPa was reached, corresponding to a strain of about 0.7%. From then on, the length of the fibre was kept constant. After the stress had during 3 hr been allowed to relax, the solvent was heated at a rate of 1° per min up to the preselected temperature. This temperature was maintained within _+0.2°.
157
J.C.M. TORFS,G. O. R. ALBERDAVAN EKENSTEINand A. J. PENNINGS
158
shrinkage(%) 15 10
-5 -10
'
100
110
=T(°C) 120
Fig. 1. Percentage of shrinkage vs temperature for a fibre heated in p-xylene in a 4 mm wide glass tube at a rate of 0.5°/min up to 100° and then at 0.25°/min. O: Upper curve: no stress. O: Lower curve: a stress of 0.4 MPa. The percentage of shrinkage is defined as A//lo x 100, where Al is the decrease of the length with respect to lo, the length at room temperature. Differential scanning calorimetry was carried out by means of a Perkin-Elmer DSC-2. The thermograms were obtained from a polyethylene fibre submerged as 4% (w/w) in p-xylene with the antioxidant dibutyl-p-cresol in sealed stainless steel pans [1 l]. The fibre was either kept at constant length by winding it on an iron wire, or it was free to shrink, because it was cut into 3 mm long pieces. Corrections for thermal lag were carried out according to the DSC manual, and indium was used for temperature calibration. RESULTS First of all, we examined the possibility of applying the equilibrium dissolution temperature of polyethylene in p-xylene of 118.6 ° [4, 5] to the fibres. In these experiments a fibre, to which no stress was applied and which was kept in a glass tube, was dissolved. When heated, such a fibre became transparent at 90°; at higher temperature, it shrank by slightly curling up to 16% at 119.4 ° as shown in Fig. 1 (solid circles). At 119.5 °, the fibre broke at several places and dissolved completely within a few minutes. These experiments proved that the unloaded fibre dissolves very dose to the equilibrium dissolution temperature of polyethylene in p-xylene of 118.6 °. When a fibre was submerged rapidly in p-xylene at 124 °, it dissolved visually within one minute, indicating the high rate of dissolution. The dissolution of unloaded fibres was also investigated by DSC of pieces of fibre, short enough to be able to retract easily. Therefore, pieces of fibre, 3 m m long, were heated in p-xylene. The ends of the endotherms were considered to be the dissolution temperatures. At heating-rates of 0.31, 0.62 and 5.0 ° m i n - 1 , the dissolution temperatures were 121.5, 121.7 and 124.0 ° respectively. The lowest value is comparable to the value of 119.4 ° obtained above at a rate of 0.25 ° per min. Simultaneous with dissolution, shrinking occurs (Fig. 1) due to the adoption of the molecules of a
more coiled conformation, first of all in the most defect regions. This seems to suggest that prevention of this coiling by applying a large enough stress to the fibre would reduce its solubility. When a small stress of 0.4 MPa was applied to a fibre, it dissolved in an essentially similar way. It started shrinking at about 100 ° (Fig. 1, open circles). At 119.0 ° it became transparent and elongated, and at 119.5 ° it broke and dissolved within a few minutes. It is remarkable that so small a stress (1/8000 of the tensile strength of 3 . 2 G P a at room temperature) is able to elongate the fibre. Elongation began as soon as the fibre, probably due to swelling, began to grow transparent. This indicates that the mobility of the long polymer molecules is very high in the swollen fibre, probably because crystallinity has disappeared. Dissolution experiments involving considerable stress were carried out with a bundle of four fibres. They were strained to obtain the initial stress and then kept at constant length. After 3 hr of stress relaxation in order to achieve equilibrium, the bundle was heated up to the preselected temperature, which was kept constant. The stress on the fibre from the moment the temperature was constant, is plotted in Fig. 2 for experiments carried out at various temperatures. (Stress was plotted instead of relaxation modulus because preliminary experiments indicated it to be the more fundamental parameter. This point will be further investigated in future.) At temperatures up to 132 °, the experiments revealed that the stress decreased most rapidly during heating and the first hours after temperature equilibrium had been established. Stress relaxation probably occurs because of taut tie molecules---which bridge disordered regions---slipping out of the crystaUites. The same process takes place in the case of the annealing of cold drawn fibrous material [12]. A few hours after reaching temperature equilibrium, the stress relaxed more slowly, and seemed to approach a limiting value. Thus the fibre seemed not to dissolve, and could still be seen after two days, even at 132.0 ° i.e. 13° above the equilibrium dissolution temperature. Experiments at 130.8 ° were carried out during 6 days and even then the fibre did not dissolve completely. Partial dissolution was investigated by weighing the fibres before and after the experiments. It was found in an experiment carried out at 130.8 ° that only (7 + 4)% of the fibre had dissolved after 3 days. This loss of weight might be the result of the existence of a separate fraction (e.g. a skin) dissolving more easily than the main part of the fibre; the possibility of very gradual dissolution however, cannot be rejected as a possible explanation at present. These results indicate that, up to 132 °, the fibre does not dissolve or does so only extremely slowly. At the higher temperature of 136.5 °, the stress relaxed fully within 7 hr. Then complete dissolution followed within a few minutes, indicating again that the stability of the fibre results from the stress. At temperatures between 132 and 136 °, stress relaxation takes place more slowly and is irregular because of differences in stress within the bundle of fibres, causing different dissolution times. Stress relaxation, followed by dissolution, still takes hours even at 136.5 °. Therefore a quickly heated
Dissolution of extended chain PE fibres 50
159
STRESS (MPa)
z.O
1
3(3
.
•
i
116 5°CL 1250 °C
1308 °C 132 O°C
lO
o
0
500
1000
1500
1 2000
TIME(ram) I 2500
Fig. 2. Stress plotted vs time for a polyethylene fibre in p-xylene at the indicated temperatures. The equilibrium dissolution temperature is 118.6C. The fibre was previously heated at a rate of l:,/min. I Pa = 10dyn/cm 2 = 1.45 x 10 4 p s i = 1.02 x 10-Skgf/cm 2 = 1.13 x 10-ag/denier(ifd= I).
sample should dissolve at a considerably higher temperature; this effect could be investigated by DSC. In these experiments, contraction of the fibre was prevented by winding it on an iron wire. Heating at 5 ° per rain resulted in an endotherm with a maximum at 150°. This indicated that the fibre was partially crystalline up to above 150 °, i.e. far above the equilibrium dissolution temperature of 118.6°!
DISCUSSION
The dissolution experiments most remarkably show the rapid dissolution of the free fibre even at 119.5 ° whereas, when considerable stress is being exerted, the fibre can withstand 132 ° for at least 40 hr. At 132 ° the temperature is higher than the equilibrium dissolution temperature; therefore, the free energy of dissolved random coils is lower than that of molecules in the fibre (this effect is even larger, because the free energy of the fibre has increased by the strain energy). Thus the fibre is metastable and its apparent thermal stability arises from a very low or zero dissolution rate, for kinetic reasons. This low rate could be ascribed to the stress, but one must also consider other possibilities. Slow dissolution might be due to the well-known decrease of rate of dissolution with increasing crystallinity [13, 14] and with increasing molecular weight, caused by the formation of entanglements (these cause an enormous increase in the thickness of the surface gel-layer through which the diffusion of solvent has to take place [-13, 15, 16]). The fibres investigated in this study are highly crystalline (about 8._0%)and consist of high molecular weight polymer (My, = 4 x 106). But there are indications that the dissolution of the fibres is not always so slow. 1. The experiments in which no stress was exerted, and which were carried out with essentially the same materials as the experiments with stress, revealed a
visual dissolution of the fibre within a few minutes at 119.5 °, and within 1 min at 124 :. 2. Differential scanning calorimetry of unconstrained 3 mm pieces of fibre submerged in p-xylene revealed that, at the scan speed of 5 ° per mira the final melting temperature is only 2.5 ° higher than at the lowest rate (0.31 ° per rain). Hence a dissolution time of only 30 sec is estimated. Therefore, the extended chain fibres do not have an intrinsically low rate of dissolution: they dissolve slowly due to the stress. In previous work [17] it was shown that the fibres consist of parallel microfibrils, themselves sequences of alternating crystallites and amorphous regions [12, 17, 18]. The thermal stability of the stressed fibres prompts two further conclusions concerning the arrangement of their chain molecules. (i) The surface of the microfibrils does not act as a less stable region, where dissolution can start due to peeling the molecules off one by one. This requires the chains on the surface to be firmly anchored inside the crystalline blocks. In other words: "The chains traverse the fibril in lateral direction on account of topological chain defects such as intertwinings, twist disclinations [19], jogs [20], loops and entanglements". (ii) The microfibrils do not break-up in the amorphous regions, in which the segmental mobility must be high due to the liquid state and to swelling. "Therefore subsequent crystalline regions have to be connected by tie molecules" bridging the amorphous zone. This conclusion is in agreement with the concept of "continuous crystal" introduced by Porter [21], because the tie molecules between crystalline regions may crystallize near room temperature. These structural features can be depicted qualitatively in a schematic figure (Fig. 3 [22]), which shows a stressed fibre at about 130 °. The existence of topological defects is not surprising in view of the mechanism of formation of these fibres in the surface growth technique [1,2]. This technique is based on the stretching of an entangled
160
J. C. M. TORFS, G. O. R. ALBERDA VAN EKENSTEIN and A. J. PENNINGS
Fig. 3. Schematic representation of molecules in a stressed fibril submerged in a solvent at a few degrees above equilibrium dissolution temperature. Extended chain crystals are partially interrupted by disordered regions [bridged by tie molecules (a)], where the concentration of defects is higher and where cilia (b) emanate the fibril. The defects depicted are trapped entanglements (c), intertwinings (d), jogs 1-20] (e), and chain ends (f). network, followed by rapid crystallization. Hence, entangiements and intertwinings occurring in the solution are trapped in the fibril. The question now is: why do not the amorphous regions grow at the expense of the crystals? This would transform the fibril into a wholly amorphous state in which entanglements can slip off, and the structure disintegrate. The answer is that the dissolved tie chains (which are in equilibrium with the crystallites) have very low entropy due to the stress restricting their number of possible conformations [23, 24]. Thus transition of a crystalline segment to the stressed amorphous state is unfavourable. Hence, dissolution of the metastable fibres does not take place, because the fibres would have to dissolve via a path involving a very high free energy barrier: a large part of a crystallite would have to swell for a moment, in such a way that the entangled chains could slip through and release, and dissolution could start at the crystallite surface. Something else adding to the free energy barrier is the relatively high polymer concentration in the dissolved parts of the fibril. Thereby, the lowering of the melting point of the pure polymer by the solvent [4,25] from 145.5° to about 118.6° [4, 5] in very dilute solutions, is not complete in the vicinity of the melting crystal face. This behaviour is distinct from the dissolution of lamellae. Then the local polymer accumulation will spread by the unfolding of the chains and by the diffusion of entirely dissolved molecules.
At this stage we can conclude that the fibre structure strongly resembles a chemically cross-linked network. Due to the topological defects, the structure amounts to an entanglement network. The entanglements cannot vanish because the chains cannot slip out of the crystallites. Therefore the theory as formulated by Gee and Flory [26-28] for chemically crosslinked networks is applicable. From this theory it can be deduced that, by application of a macroscopic stress, the melting temperature of extended chain .crystallites rises. The stress is transduced to all crystallites by the tie molecules. The behaviour of the strained fibres is different above 132 °. Stress relaxes completely and then dissolution occurs as in the case of free fibres. Stress-relaxation indicates that the mobility in the crystals is so high that tie molecules can slip away. It should be noted that this limit of 132 ° closely approaches the limit for fibrous growth by means of the surface growth technique, which uses the same solvent. When the cross-sections of the fibres are plotted vs crystallization temperature and an extrapolation towards zero cross-section is made, the achieved temperature is about 130° [29]. This suggests that growth would be impossible at 130°. The temperature-limit thus obtained for the surface growth technique might very well arise from too rapid slippage of chains past each other, resulting in rapid creep. The stress relaxation temperature is also remarkably close to 130 °, where the lamellar thickness of melt-crystallized or long time annealed solutiongrown polyethylene increases by a factor of 2-5 [30]. This rather sudden increase in lamellar thickness suggests a higher chain mobility, dependent on temperature but not on a solvent (since the experiment was carried out without solvent). "The consistency of the three temperatures (stressdissolution: 132°; surface growth: 130°; annealing: 130°: suggests a sudden higher mobility for polyethylene chains in the crystalline phase." This is probably due to thermally generated defects, already significant at 110 ° [31]. Finally it should be remarked that the dissolution temperature of 119.5° of free fibres must be considered high. As a result of the high free energy of the crystal surface and defects, a dissolution temperature below the equilibrium value of 118.6 ° would be expected. Dissolution of polyethylene has indeed been observed within the range of 110-111 ° for high molecular weight lamellar crystals [4] and for high pressure crystallized extended chain crystals [32-34]. However, the behaviour of stacked extended chfiin crystallites connected by tie molecules (Fig. 3) must be different e.g. from re-entering chains in lamellae. In the former case dissolution takes place at a higher temperature because the entropy of the dissolving tie chains is relatively low, mainly because of the limitations on the number of conformations. Limitations result from the volume occupied by the crystallites and from the fixation of the tie chains in the crystallites [23, 24] they connect. Work is now in progress on other materials and on such aspects as the influence of initial strain and solvent quality on the dissolution process of these extended chain fibres.
Dissolution of extended chain PE fibres CONCLUSIONS The fibres produced by surface growth [1-3], containing extended chain crystallites, dissolve at a higher temperature than e.g. lamellar crystals. This effect is most pronounced when the fibres are subjected to a stress. Then the apparent dissolution temperature rises to 1 3 above equilibrium dissolution temperature (in p-xylene). The thermal stability is indicative of: 1. Tie molecules bridging the a m o r p h o u s regions between crystallites in a microfibril. 2. Topological defects such as entanglements. Thus the structure a m o u n t s to a network of entanglements which are trapped by crystallites. Crystallites do not dissolve, because the entropy of the stressed tie chains in the dissolved state is much lower than the entropy of free molecules in solution. Above 132 the fibres dissolve, because the applied stress relaxes, due to high chain mobility in the polyethylene crystals.
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