Studies of orientation phenomena in crystallizing polymers

JOURNAL OF POLYMER SCIENCE VOL. XXI, PAGES 363-379 (1956)

Studies of Orientation Phenomena in Crystallizing Polymers

A. KELLER, * Imperial Chemical Industries Limited, Dyestufs Division, Manchester 9,. England

INTRODUCTION

This paper records some anomalous orientation efiects observed in the course of our work on various crystallizing polymers. An interpretation of these effects will be attempted whenever possible. The paper consists of two parts: (I) extrusion, (11) drawing. By extrusion is meant the formation of a filament directly from the melt, and by drawing the further elongation of the solid crystalline extruded material. These terms as used here are quite general and they do not necessarily refer to any of the actual technical processes of extrusion and drawing, as adopted in the manufacture of the various fibers. The object of the present investigations was to detect anomalous structural features, such as can occur when fibers are made. In a given material these structures can be more or less pronounced, and can occur more or less frequently depending on the particulars of the orientation process.

PART I. EXTRUSION

In the course of fiber manufacture the molten polymer is first extruded through holes of required diameters, and the solidifying filaments are wound up on rollers. The product obtained in this way is then drawn further in order to arrive at the final yarn. Usually no crystallite orientation is found in the extruded filament. Occasionally observed weak birefringence effects have been interpreted as being due to some incipient orientation in the amorphous portions of the molecules. In the case of high wind-up speeds orientation can occur and according to the earlier findings this orientation is an imperfect version of the final fiber orientation, as if incomplete drawing had taken place (e.g., reference 1).

It appears from our experiments that extrusion can produce orientation within the crystalline parts, even if the material is not being wound up at all, consequently when under slight tension only. This orientation was found to be of a new type, which does not correspond to a simple parallel ordering of the chains as would be expected if some incipient “drawing”

* Present address: H. H. Wills Physical Laboratory, University of Bristol, Bristol, 8. England.

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364 A. KELLER

took place during extrusion, but appears to be characteristic of the extrusion process itself (extrusion being understood as fiber formation from the melt without drawing).

Experimental

Most of the extrusion experiments were carried out by pressing the molten polymer through dies of 0.5-2 mm. diameter. This die was at the lower end of the heated and vertically held tubular container. The extruded material was then allowed to drip freely while it solidified in the air forming a continuous filament. The length of the free drop was about 2 feet, which, under the circumstances, was sufficient for crystallization to set in, except for the more slowly crystallizing polymers like Terylene (polyethylene terephthalate) and nylon 6,lO (polyhexamethylene sebacamide) . The speed of the flow was controlled as far as possible by the pressure exerted on a plunger, while it also depended on the temperature of the melt and the diameter of the die. This method was used mainly with the polyamides and occasionally with polythene. Some other simple manual methods of extrusion used with polythene and Terylene will be mentioned when de- scribing the individual experiments.

Polyam ides In the polyamides extruded in the above way, an orientation, as shown

by Figure 1, could be very frequently observed. The characteristic feature

Fig. 1. Nylon 6,6, extruded.

of this orientation is the meridional position of the maxima in the inner ring (this is the 100 in nylon 6,6 and 6,10 (polyhexamethylene adipamide and sebacamide) and the 200 in nylon 6 (polycaproamide)), and the equa- torial maxima in the second strong ring of somewhat larger diameter (this is the 010 in nylon 6,6 and 6,lO and the 002 in nylon 6).2s3 This latter set of planes contains the hydrogen bonds. Such orientations where a hkO reflection is on the meridian (instead of being on the equator as in the usual fiber photographs, used in the indexing of the reflections) are termed “perpendicular,” as they correspond to the direction of the molecules, being

ORIENTATION PHENOMENA IN CRYSTALLIZING POLYMERS 365

perpendicular to the fiber axis, instead of being parallel to it, as it is usual in fibers. The strength of this type of orientation (we mean that of the above perpendicular orientation, the only type of orientation observable in these extruded samples) varied from sample to sample.

The orientation in question was first observed with nylon 6,6, where it occurred most readily; in fact it was difficult to obtain a sample which did not show i t all, however weakly. This orientation was the more pronounced the lower the temperature of fusion and the thinner the extruded filament. The most pronounced effects were observed when fusion temperature was about 265-268OC. and the filaments were about 0.1 mm. thick, Subsequent annealing a t elevated temperatures, or in boiling water, improved the orientation (perpendicular orientation is meant all the time) and also sharpened the definition of the x-ray pattern.

Nylon 6 behaved very similarly to nylon 6,6. The perpendicular orientation was not always as pronounced but i t could be brought out very strongly by subsequent annealing. Material which was received as “un- washed,” i.e., which contained more monomer, showed the orientation effects in question much more readily.

Nylon 6,lO showed much weaker perpendicular orientation, noticeable only after subsequent annealing. Obviously this is due to the fact that this polymer crystallizes more slowly, so that only little crystallization could take place during the short time while the filament solidified. For the same reason this technique of extrusion was still less suitable for similar experimentation with Terylene, a polymer which is completely quenchable.

The birefringence of all these polyamides was always positive. No definite structural detail could be seen when examined.with the microscope.

It is of interest t o compare Figure 1 with Figure 2. Figure 2 was given by a section cut from a block of nylon 6,6 moulding. For this purpose sec- tions were cut from molded blocks, and it was found that these contained flow lines and corresponding “row-type” ~pherulites.~ In some moldings these rows were parallel over the entire cection (Fig. 3). Figure 2 was ob-

Fig. 2. Nylon 6,6, moulded, containing parallel rows as in Figure 3.

366 A. KELIXR

tained from such a section with the direction of the rows being vertical. It is seen that Figure 2 corresponds to Figure 1 turned through a right angle. The same effect as in Figure 2 was observed in some earlier experiments with monofils. Here undrawn filaments of nylon 6,6 were tied on a hori- zontal wire, and one end of the latter was heated with a flame. The heating of the wire could be regulated so that the filament just softened where in contact with the wire, and started to pull out a t these localities. The pulling out was regulated with small weights hung on the free end of the mono- fil. In this way two types of perpendicular orientation could be produced, giving x-ray patterns as Figures 1 and 2. The former was positively, the latter negatively, birefringent. The latter effect was very rare and occurred only along very short lengths (1 mm.). These latter results were briefly mentioned in reference 4, and were interpreted as alignments of the two types (positively and negatively birefringent) of spherulite radii (helices). This interpretation proved to be partly erroneous. Subsequent examination revealed that the negatively birefringent samples had the row structure as Figure 3.

Fig. 3. Photomicrograph of a section of moulded nylon 6,6. Crossed Nicols X 135.

Polythene

Polythene of low grade number (grade 7 and below: high molecular weight above 28,000 6, did not flow out readily. The molten material appeared to be quite elastic and could be pressed out only as one coherent column without flow. When the hanging molten filament was induced to flow by manual pulling the OW" orientation5 was produced. With polythenes of high grade number flow did set in spontaneously, but the resulting filament showed no orientation. Some orientation, however, could be achieved with material of very high grade number (7000; molecu-

ORIENTATION PHENOMENA I N CRYSTALLIZING POLYMERS 367

lar weight low, 5300 6, when the melt was pulled out from a beaker with a glass rod. The orientation was excessively faint and it could only be detected in photographs so weakly exposed that the reflection under examination just about became visible. The three principal reflections (110, 200, 020) are all of difTerent strength; consequently each of them was examined on a separate photograph of suitable exposure. Independent examination of these three reflections unambiguously indicated a perpendicular orientation as sketched in Figure 4 (e and a axes perpendicular, b parallel, to the

Fig. 4. Polythene grade 7000, extruded (sketch).

length direction). Because of the extreme weakness of the effect the sample is still to be regarded as essentially unoriented. However, the mere fact that this structure does exist in polythene is significant as in principle it completes the analogy between polythene and the polyamides (see Discus- sion).

Terylene Molten Terylene could not be made to flow slowly enough to allow

crystallization to set in before the treacly material broke. (One isolated example in which this operation was partly successful is given in Part 11.) For this reason, amorphous strips cut from quenched ribbons were placed on a temperature gradient bar between temperatures of 210 and 25OoC., when they softened instantaneously before they had a chance to crystallize while heating up. The soft strips were placed over a different portion of

368 A. KELLER

the temperature gradient bar at the temperature of about 180°C., and were pulled slowly by hand until the material crystallized. The x-ray patterns were usually poorly defined and contained a large proportion of amorphous material, as seen from Figure 5 . The pattern in Figure 5 corresponds to a

Fig. 5. Terylene undried, “extruded.”

perpendicular orientation as reflected by the meridional position of the maximum in the 100 ring (outermost rings). The maxima of the 010, O i l reflections (the innermost rings) are clearly in the region of the equator. (For structure see reference 7.)

Interpretation and Discussion

All the extruded samples, described above, showed perpendicular orientation, whenever orientation was noticeable. Perpendicular orientations have been described previously by the author and also by others (see 4 and 5 and references there). This was found to be the type of orientation which forms spontaneously d iring crystallization from the melt. In the absence of external orientation effects it only extends over microscopic regions (spherulites), while under the influence of flow it can give perpendicular character to the orientation of the whole macroscopic sample (rows6). Further, perpendicular orientations were found also in the drawn samples which were relaxed at elevated temperature^.^^^ This latter type of perpendicular orientation, however, differs from that produced by spontaneous crystallization. In principle several perpendicular orientations can exist. The direction of the molecular chains is always perpendicular to the axis of rotation; however, the crystallites could have any of the possible positions around this direction. Except for Terylene the position of the unit cell around this direction is defined by one of the other two principal axes being also perpendicular to the axis of rotation, and in the same polymer this is a different principal axis in the spherulitic and in the relaxed samples. Ex- ceptionally the perpendicular orientation characteristic of the relaxed samples can also appear in the spherulites (negative spherulite of nylon 6,10)8~9 (see relaxation experiments below in Part 11).


The existence of perpendicular orientation is unexpected in high polymers. Various considerations have led the author to suggest that this type of orientation is caused by a closely coiled arrangement of fibrils. One such consideration was the apparent individuality of fibrous units along the radii of the spherulites. This feature is inconsistent with the earlier picture (e.g. , references 10 and l l ) , which visualizes these crystalline fibrous units as being connected by the chain molecules so that one chain passes through several of them, the radiating crystslline units themselves being formed through the alignment of the molecules over a part of their lengths. Ac- cordingly the fibrils would be formed through intermolecular forces, while the different fibrils would be connected with each other by valence forces. It has been pointed out recently by Reding and Brown,12 and is also borne out by the experimental material of this present work, that such a model is untenable, while even apart from any other consideration the helical model is a t least consistent with all the observed experimental facts. Therefore, the radiating crystalline units building the spherulites will be frequently referred to as helices throughout this paper.

It was shown in our work on polythene that two different types of perpendicular orientations can give very similar x-ray patterms The present findings on polyamides will be considered in the light of these earlier results on polythene. A comparison of Figure 1 with that of Figure 1 in reference 9 reveals that the extruded filament gives the same diffraction pattern as a fibril lying along the radius of a positive spherulite. This suggests that the extruded filaments consist of such fibrils aligned in the extrusion direction (structure a). However, a second alternative would give the same diffraction pattern. Here the fibrous units which build the negative spherulites would be the structural units, but oriented perpendicular to the filament direction, an arrangement which would be realized by the row structure produced by negative spherulites (structure b) . The analogy with polythene would suggest a row-type structure, i.e., structure b, but on the other hand negative spherulites are extremely rare in the unoriented material, a consideration which would favor structure a. Most of the methods used to distinguish between the two alternatives in our earlier work on polythene5 could not be applied to the present problem. There is no con- clusive microscopic detail; also the sign of birefringence is expected to be positive in both cases. Definite evidence, however, could be obtained from the 002 reflection. In structure a, the maxima are expected to be a t an angle of loo, but in structure b, a t an angle of 40' as measured from the equator. It is seen from Figure 1 that the first consideration is satisfied, accordingly we are dealing with structure a. As regards Figure 2, we know that the pattern is given by the usual row structure, i.e., by the helices which normally lie along the radii of the positive spherulites, but here they are all oriented perpendicular to the flow direction. Accordingly Figure 1 is given by the same structural units as Figure 2, only turned through a right angle as suggested by the first inspection of the photographs.

In view of what has been said the interpretation of Figure 4, given by

370 A. KELLER

polythene, offers no difficulty. The pattern is identical with that of Figure 3 in reference 9, which means that the structure of the filament is again identical with that along the radius of a spherulite. There is no reason to suspect the presence of an alternative row structure, because we have no knowledge of any other type of spherulite as in the case of polyamides. Consequently Figure 4 is related to Figure f in plate I of reference 5 as Figure 1 to Figure 2 in the present paper.

With Terylene (Fig. 5) the possibility of the row structure does not arise. On the basis of Figures 4 and 5, both in reference 9, this would require the innermost strong reflections (011, 010) to be near the meridian, which is clearly not the case. It is seen that Figure 5 comes nearest to Figure 5 of reference 9, which means that the orientation in the strip is very similar to that along the radius of a negative spherulite occurring in undried samples.8 The fact that the structure in the extruded strips corresponds more nearly to the negative than to the positive spherulites is to be expected, because the original strips were undried. The same extrusion experiment could not be carried out with dried samples, as dried samples cannot be obtained amorphous, because of crystallization setting in a t the elevated temperatures needed for drying.

It is seen that Figure 5 and Figure 5 in reference 9 are not quite identical. In the former the 100 plane is tilted through a larger angle, apparently through about 90' with respect to the fiber axis. This suggests that the perpendicular orientation (or nearly such) is not uniquely defined, but it can be different in different samples. This is borne out by some observations made on unoriented, partly dried samples of Terylene (the mention of which was omitted in the earlier paper on spherulites). Here spherulites could be seen which were nonbirefringent, but still visible with unpolarized light. Further, in the same samples positive and negative spherulites could be seen together, both of which appeared in continuous degrees of brightness between the maximum brightness and invisibility when examined between crossed Nicols. As the sizes were about the same, this is obviously due to a continuous range of birefringence between a maximum positive and maximum negative value. Consequently one might expect orientations intermediate between Figures 4 and 5 (both in reference 9) or such positive and negative orientations which differ more from each other than do Figures 4 and 5 (both in reference 9).

Individual rows could be often observed microscopically in unoriented samples of Terylene, but we could not obtain samples consisting entirely of such rows. The possibility of row orientation within the whole sample, however, is not excluded.

These experiments all bear out the striking conclusion that the same operation, i.e., extrusion of the crystallizing melt, can produce two types of orientation within the same polymer (with the apparent exception of Terylene and even there the experimental possibilities are not exhausted)- both arising from the same morphological units (fibrils, helices) with the difference that in one case these units are parallel, while in the other they

ORIENT-4TION PHENOMENA IN CRYSTALLIZING POLYMERS 371

are perpendicular, to the extrusion direction. The first case corresponds to an alignment of these helices themselves, the second to an alignment of larger aggregates built by them. The sorting out of these factors was possible only with the help of the morpholical approach, which considers polymers as consisting of units which themselves are crystal aggregates.

It could be asked, why do row structures form most readily in polythene, less frequently in polyamides, and only in isolated instances and over small localities in Terylene? Further, why do fibrous units align readily in polyamides while only to a very slight extent, and even then only under special conditions, in polythene? No explanation is attempted a t this stage. Purely empirically i t appears that the formation of row itructures is favored by high viscosity and by a low value of the ratio (rate of flow)/(rate of crystallization).

The finding that some of the extruded samples have a structure which is identical with that along the radius of a spherulite in the same unoriented material has an interesting implication. The filament having such a structure can be considered as the radius of a giant spherulite.

Fig. 6. Nylon 6,lOcrystallized at 160 "C., Fig. 7. Nylon 6,6 drawn at about drawn 60%. 130 "C.

PART 11. DRAWING

Experimental

Polyam ides Polyamide strips (nylon 6,6 6,10, and 6) were drawn slowly in a frame

up to an elongation of about 60%, the elongation still being homogeneous (no necking). The strips were elastic within this region of elongation, and recovered nearly completely when the stress was removed. Figure 6 illustrates an x-ray pattern obtained in the stretched state. It is seen that the orientation is perpendicular and very similar to Figure 1 and to Figure

372 h. KELLER

Fig. 8. Nylon 6, extruded, annealed at 190°C.

1 in reference 9. The sign of birefringence was positive, and, in some samples which contained large spherulites in the unstretched state elongated, distorted spherulites could be seen under the polarizing microscope when the draw direction was nearly parallel to the vibration direction of either the analyzer or the polarizer.

A t room temperature 'further elongation was no longer possible without necking. This took place after an apparently high yield point had been reached (as judged from manual pulling). The orientation always corresponds to the usual fiber orientation.'% When the material was strongly crystalline in the unoriented state, drawing had to be carried out slowly and cautiously to avoid rupture. In this latter case the definition of the diffraction pattern deteriorated as a result of drawing in nylon 6,lO (Fig. 9 in reference 13). Nylon 6 behaved differently in that the definition of the dzraction rings improved with drawing (Fig. 16, reference 13).

Fig. 9. Nylon 6, extruded, drawn 70% Fig. 10. Nylon 6, extruded, drawn 280% at 190" C. at 190°C.


When these polyamides were drawn at higher temperaturs, drawing could be achieved continuously without necking (a well-known fact). For moderate extensions the resulting orientation was strongly planar, as seen from Figure 7, where the inner strong ring extends over most of the circle while the outer one is concentrated in the region of the equator (hydrogen- bonded planes orient first). Occasionally selective orientation could be detected (maxima along the inner ring between equator and meridian).

Particularly interesting were some of the experiments carried out with nylon 6, previously extruded in the way described in Part I. The starting material had the usual perpendicular orientation. This material was drawn at 190°C. Figure 8 is the diffraction pattern given by the starting material, held a t 19OOC.' for the same length of time as the drawn samples giving Figures 9 and 10. The sample giving Figure 9 was drawn to about 70% elongation without necking, the one giving Figure 10 was drawn to 280Y0 elongation, during which it necked. As seen, the orientation is perpendicular in spite of drawing, in fact the perpendicular character of the orientation has even improved. Only after the most extensive elongation (600%) did the conventional fiber pattern appear with the molecules parallel to the draw direction.

It was found that quenched and drawn nylon 6,lO samples could be relaxed beyond 200Y0 when placed suddenly on a heated plate of 200-210OC. The heating up had to be sudden to avoid intensive crystallization before relaxation could occur. Strongest relaxation effects could be achieved by moving a drawn filament rapidly through the flame of a burner. This simple experiment was particularly instructive because it produced a relaxation gradient containing all stages of relaxation from the unchanged drawn part to the completely relaxed probably molten and recrystallized part, within the same sample. Relaxation proceeded through selective orientations, where the crystallites tilt around their a axes, the orientation of the 100 planes being undected. Figures 11 and 12 show two stages of the series; it is seen that in Figure 12 perpendicular orientation is approached.

Fig. 11. Nylon 6,10, moderately Fig. 12. Nylon 6,10, strongly re- relaxed. laxed.

374 A. KELJXR

Polyfhene Drawing and relaxation studies have been described in reference 5.

Terylene No systematic work was carried out with Terylene here; only some un-

usual observations will be mentioned. Amorphous strips of Terylene were drawn in hot air (in a heated metal

box with mica windows) a t 14OoC., slowly so as to allow crystallization to set in. The samples gave dzraction patterns as Figures 13 and 14.

Dried crystalline material was fused over a flame and drawn a t 180°C. (in a similar arrangement as above) while crystallization sets in. This experiment would really belong to Part I, being more like an extrusion; i t is listed here because of the nature of the resulting diffraction pattern shown by Figure 15 (see Discussion).

Amorphous Terylene was drawn above 200"C., and the resulting filament was then relaxed by moving it cuickly through a flame. The diffraction pattern is shown in Figure 16.

Figs. 13-16. Terylene drawn hot. For details see text.


Interpretation and Discussion

In the first stage of the drawing of polyamides (Fig. 6) the resulting orientation corresponds to that along the radius of a positive spherulite. Such an orientation would arise if the helices which we visualize as forming the spherulites are aligned with respect to the draw direction, which is in agreement with the interpretation of drawing put forward earlier.4 The high elasticity indicates that this alignment is achieved by a reversible distortion of the spherulites, without a major disruption of these units. This is also borne out by our microscopic observations. Elongated distorted spherulites have also been reported by earlier author^.'^^^^ Our findings here demonstrate that this distortion is reflected by the over-all orientation of the crystallites in the sample as a whole.

(It is of interest to note that this morphological interpretation has an early precedent. As far back as 1898 Ambronn noticed that gutta-percha became negatively birefringent during the first stages of elongation (reference 16, as quoted by Wachtler)." He interpreted this behavior as resulting from an alignment of fibrillar units which form the negatively birefringent gutta-percha spherulites, the existence of which was apparently known at that time.)

The similarity of Figures 1 and 6 raises the question whether the extruded samples discussed in Part I consist of such spherulites, strongly distorted while forming within the flowing melt, or whether they consist simply of more or less unrelated helices growing in the flow direction instead of growing radially outward from a center. The latter case appears to be the likely one, because there is no indication of a force tending to restore the original shape of the spherulite as in samples corresponding to Figure 6. In the extruded samples the definition of the perpendicular orientation did not decrease even when the samples were heated subsequently to the melting range; on the contrary, it improved, presumably due to improved paral- lelization of the helices, and to additional crystallization and also recrystal- lization, which followed the existing pattern of orientation already present. Consequently the slightly drawn and the extruded samples corresponding to Figures 6 and 1, respectively, are not identical in spite of the similarity of the orientation, as the former is the result of spontaneous crystallization, and accordingly appears to be free of macroscopic strains.

When drawn at room temperature, the sample resists further elongation beyond the stage shown by Figure 6, until necking occurs at much higher stresses. A t this stage the orientation changes discontinuously from perpendicular orientation to the final fiber orientation. If drawing is carried out a t elevated temperatures there can be a more or less continuous distri- bution of orientation between the two extremes within the sample. It is of interest to compare this behavior with that of p~lythene.~ In polyethylene drawing produced selective orientation, even a t the first stages. The final fiber orientation was approached gradually on further drawing. This was interpreted by an alignment and gradual pulling out o'f helices forming the

376 A. KELLER

spherulites. The pulling out itself would occur so readily that i t is not possible to avoid i t completely, which would account for the fact that perpendicular orientation cannot be produced by drawing. The different behavior of polyamides must be due to the stronger cohesion within the units (helices) which form the radii of the spherulites. This is borne out by the fact that these can be aligned without altering their internal structure (perpendicular orientation), and that their final pulling out needs a greater force. In fact they cannot be pulled out smoothly a t room temperature, and the high yield point indicates that some internal rupture is taking place.

This last point is reflected by the decrease of crystallinity and probably also by the decrease of the crystallite size in the fully drawn material.I3 Since the force a t the yield point is greater than needed to establish a given selective orientation, once rupture occurs in the crystalline regions the final orientation is established in one stage. A t elevated temperatures the yield point is lower, consequently intermediate orientation could be a t least partly realized as reflected by the strong planar orientation and occasional traces of selective orientation. The increase in the crystallinity of nylon 6 beyond the yield point13 is a factor probably additional to the above picture and probably represents additional strain-induced crystallization. A more detailed examination of the stress-strain properties was beyond the scope of this work.

The experiments with extruded nylon 6 filaments (Figs. 8-10) are particularly illuminating. The observed behavior can be readily interpreted if we visualize the extruded starting material as consisting of an arrangement of parallel fibrillar units. Under our conditions of drawing the paral- lelism of these units would improve, and under increasing force they could glide past each other without any change in their internal structure. This illustrates the underlying hypothesis, that the forces between the gliding units are weaker than those within them. Consequently the former ones cannot be valence forces as i t is generally visualized in models of spherulite structure.10.1' These findings mean that the units in question accommo- date the whole molecule in the majority of cases a t least, and a t the same time the crystallites must be oriented so that the chains forming them should be perpendiculgr to the length of the fibril. These, to some extent incompatible conclusions, can be readily reconciled by attributing a helical morphology to the fibrous units (see above and references 4 and 9).

The results of the relaxation experiments with nylon 6,lO are in line with earlier findings. As opposed to drawing when the tilt is around the b axis (after this axis has become perpendicular to the draw dircction) relaxation occurs by a tilting of the crystallites around the a axis. The traces of this effect have been noticed first by Wilbourn and Smith (quoted in 4). Our present observations supplement these first findings by producing relaxa- tions as high as 2007& (previously only slight retractions were realized), giving all stages of pronounced selective orientation. This completes the analogy between the drawing relaxation behavior of polyamides and polythene.5


The problem of necking needs a brief mention. In these experiments necking was observed in connection with three dit€erent structural mecha- nisms: (a) cold-drawing of nylon 6,10, when orientation was accompanied by a decrease of crystallinity; (b) cold-drawing of nylon 6, which resulted in an increase in crystallinity; (c) hot-drawing of extruded nylon 6 (pos- sessing perpendicular orientation) involving the gliding of otherwise unchanged submicroscopic units. It is obvious that necking is a macroscopic phenomenon, not connected with any particular structural mechanism. The problem of necking has been dealt with in detail by Marshall and ThompsonLs and by Ja~ke1. l~ Among others these authors postulate a melting of the crystallites during cold-drawing. We suggest that “breaking up” of structures would be a more comprehensive term, which would allow for disruption on any structural level, including also the molecular melting as one extreme case.

The greater cohesion within the polyamide helices as compared with polythene might be explicable. It has to be remembered that the hydrogen bonded planes are parallel to the radii of the most frequently occurring positive ~pherulites.~ If we visualize the units along these radii as fibrils consisting of winding ribbons, the planes of these ribbons would contain the hydrogen bonds.9*zo Consequently these ribbons would be expected to be broad and even adjacent windings could be bridged by hydrogen bonds which would prevent the ribbon from pulling out without rupture occurring first. Such a helix will not be distinguishable from a continuous tube or column. The essential point, borne out by experiment, is that if we con- sider those molecules which form part of the units (column, tube, or helix), the majority of these molecules will be contained by one unit only along their entire length, the same unit accommodating both the crystalline and amorphous portions.

Figures 13-16 illustrate different selective orientations in Terylene. The exact nature of the tilts can be read ofl directly from a reciprocal lattice construction such as in Figure 6 of reference 9. Figures 13 and 14 show two stages of the same type of selective orientation. The crystallites tilt around the normal to the 170 plane through about 10’ and 20°, respectively.

In Figure 15 the normal to the 170 planes tilts in the plane defined by the fher axis and the normal to the 110 planes in the untilted position, the tilt itself being about 50’. This tilt direction corresponds roughly to that in the positive spherulites,g the magnitude of the tilt being intermediate between that in the fully drawn fiber and that in the positive spherulites, which according to our interpretation would correspond to an alignment of partly pulled out helices of the type forming the positive spherulites. The correspondence with the positive spherulites is consistent with the fact that the samples were dried prior to fusion (condition of the formation of positive spherulites8).

In Figure 16 the tilt direction can be defined as follows: If we visualize a type of tilt in which the normals to the 100 planes tilt in a plane defined by the fiber axis and by the original untilted normals to the same planes,

378 A. KELLER

then the tilt direction in Figure 16 is intermediate between the one just described and the one realized in Figures 13 and 14.

These examples illustrate the variety of tilt directions which can exist in Terylene. These do not necessarily correspond to principal crystallographic directions. In fact, they rarely do so. The first tilting effect found origi- nally by Daubeny, Bunn, and Brown’ corresponds to a tilt around the normal to the 230 planes. In the light of the present work, this is only one of the possible tilt directions, and is nearest to the tilt realized by Figures 13 and 14, but much smaller in magnitude.

References

(1) E. Kordes, F. Giinther, L. Buchs, and W. Galtner, Kolloid Z., 119,23 (1950). (2) C. W. Bunn and E. V. Garner, Proc. Roy. Soc., A189,39 (1947). (3) D. R. Holmes, C. W. Bunn, and D. G. Smith, J. Polymer Sci., 17, 159 (1955). (4) A. Keller, J. Polymer Sci., 11, 567 (1953). (5) A. Keller, J. Polymer Sci., 15,31 (1955). (6) J. Harris, J. Polymer Sci., 8, 353 (1952). (7) R. de P. Daubeny, C. W. Bunn, and C. J. Brown, Proc. Roy. Soc., A221, 541

(8) A. Keller, J. Polymer Sci., 17, 291 (1955). (9) A. Keller, J . Polymer Sci., 17, 351 (1955).

(1954).

(10) W. M. D. Bryant, J. Polymer Sci., 2,547 (1947). (11) Fibresfrom Synthetic Polymers, edited by R. Hill, Elsevier, 1953. (12) F. P. Reding and A. Brown, Ind. Eng. Chem., 46,1962 (1954). (13) I. Sandeman and A. Keller, J. Polymer Sci., 19, 401 (1956). (14) C. M. Langkammerer and W. F. Catlin, J. Polymer Sci., 3,305 (1948). (15) E. Jenckel and E. Klein, Kolloid 2.. 119,86 (1950). (16) H. Ambronn, Ber. d . Kgl. Stichs. Ges. d . Wiss. math.-physik, Kl . Natunu. Teil, 50,

(17) M. Wiichtler, Fortschr. d . Minerdogie, Krisl. u. Pelrogr., 12, 643 (1927). (18) I. Mtirshall, and A. B. Thompson, Proc. Roy. Soc., A221,541(1954). (19) K. Jiickel, Kolloid Z., 137, 130 (1954). (20) A. Keller, Nature, 171, 170 (1953).

643 (1898).

Synopsis

Unusual orientation effects encountered during the formation of filaments of polyamides (nylon 6,6, 6,10, and 6 (polyhexamethylene adipamide, sebacamide, and polycaproamide)), polythene, and Terylene (polyethylene terephthalate) are described and discussed. The paper is divided into two parts: ( I ) extrusion, (11) drawing. ( I ) It was found that, under certain conditions, the extrusion process can produce a characteristic perpendicular orientation of the crystallites. This type of orientation can correspond to two types of structures in the extruded samples: ( 2 ) the row structure, which consists of aggregates of partially developed spherulites forming along flow lines; (2) a new structure, in which the crystalline fibrillar units (helices in our interpretation) grow in the extrusion direction only, instead of forming the normal spherulites. Both structures occur in all the polymers investigated, but their relative frequency depends on the material and on the conditions of extrusion. (11) It was found that polyamides have a perpendicular orientation in the first stages of elongation corresponding to a strong elastic distortion of the spherulites, and further, that rupture occurs within the crystalline region during the later stages of drawing. It was shown that, under certain conditions, drawing can consist in the gliding past of crystal aggregates which themselves possess perpendicular orientation. Attention is drawn to some unusually pro-


nounced relaxation phenomena observed in polyamides and to the variety of crystallographic tilts, which can occur in Terylene. These experimental findings are discussed. In the course of this discussion the usefulness of the morphological approach in the treatment of these types of phenomena is clearly brought out.

RBsum6

Des effets d’orientation inusuels ont 6th dCcrits et discutCs en cours de formation de filaments de polyamides (nylon 6,6,6,10 et 6; polyhexamCthyleneadipamide, sebacamide et polycaproamide) de polythsne et de tCrylene (tCrCphtalate de polyCthylene). Ce manuscript est divise en deux parties: ( I ) extrusion, (11) la filature. ( I ) On a trouvC que dans certaines conditions le procCdC d’extrusion peut produire une orientation perpendiculaire caractkristique des cristallites. Ce type d’orientation peut correspondre A deux types de structure dans les Cchantillons extrudes: ( 2 ) la structure consistant en une aggregation de sphCrulites partiellement dCveloppCs et form& le long des lignes d’Ccoulement et (2) une structure nouvelle oh les unit& fibrillaires cristallines (hClices dans notre interpretation) croissent dans la direction de l’extrusion uniquement, au lieu de former des sphCrulites normaux. Ces deux structures se prksentent dans tous les polymeres CtudiCs, mais leur frCquence relative dCpend du matkriau utilisk et des conditions d’extrusion. On a trouvC que les polyamides prCsentent une orientation perpendiculaire dans les premieres Ctapes de 1’Clongation qui correspondent A une forte distorsion Clastique des sphCrulites, et ensuite que la rupture se passe dans les dernieres Ctapes de 1’Ctirement. Dans certaines conditions 1’Ctirement peut consister dans le glissement de la pihe des aggrCgats cristallins qui posddent eux-m&nes une orientation perpendiculaire. On attire l’attention sur certains phCnom8nes de rClaxa- tion exceptionnellement prononds dans les polyamides et sur la variCtC des agencements cristallographiques, qui se prCsentent dans le tCryl8ne. On discute ces resultats exp6ri- mentaux. On dCmontre en cours de cette discussion l’utilith de I’approche morpholo- gique dans 1’Ctude de phCnomenes de ce genre.

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Zusammenfassung

Ungewohnliche Orientationswirkungen, die wahrend der Bildung von Fasern von Pol yamiden (Nylon 6,6, 6,lO und 6 (Polyhexamethylenadipinamid, -sebacinamid und Polycaproamid)) Polythen, und Terylen (Polyathylenterephthalat) auftreten, werden beschrieben und diskutiert. Die Arbeit ist in zwei Teile eingeteilt: ( I ) Ausstossung; (11) Streckung. Es wurde gefunden, dass der Ausstossungsvorgang unter gewissen Bedingungen eine charakteristische senkrechte Orientation der Kristallite hervorrufen kann. Diese Art der Orientation kann mit zwei Strukturarten in den ausgestossenen Proben iibereinstimmen: ( 2 ) die Reihenstruktur, die aus Aggregaten von teilweise entwickelten Spheruliten besteht, die sich an den Fliesslinien entlang bilden, und (2) eine neue Struktur, wobei die kristallinen faserartigen Einheiten (nach unserer Ausle- gung Schneckenformen) nur in der Ausstossungsrichtung wachsen, anstatt die norrnalen Spherulite zu bilden. Diese beiden Strukturen treten in allen untersuchten Polymeren a d , aber ihre relative Hiufigkeit hangt vom Material und von den Ausstossungsbedin- gungen ab. Es wurde gefunden, dass Polyamide in den ersten Stufen der Ver- langerung eine senkrechte Orientation haben, die mit einer starken elastischen Verzer- rung der Spherulite iibereinstimmt, und weiterhin, dass in den spateren Stufen der Streckung Bruch auftritt. Es wurde gezeigt, dass Streckung unter gewissen Bedin- gungen aus dem Vorbeigleiten *on Kristalaggregaten bestehen kann, die ihrerseits senkrechte Orientation haben. Die Aufmerksamkeit wird auf einige ungewohnliche Relaxationsvorgange geleitet, die in Polyamiden beobachtet werden, und auf eine Anzahl von kristallographischen Neigungen, die in Terylen auftreten konnen. Diese experimentellen Befunde werden diskutiert. Im Laufe dieser Diskussion wird die Niitzlichkeit der morphologischen Annaherung bei der Behandlung von Vorgangen dieser Art klar hervorgebracht.

Received May 7, 1956

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Documents

Studies of orientation phenomena in crystallizing polymers