High strength elastic polypropylene

High Strength Elastic Polypropylene

ROBERT JOEL SAMUEL&* Hercules Incorporattd, Research Center, Wilmington, Delaware 19899

Synopsis

The transformation of a spun isotactic polypropylene fiber into a high strength elastic fiber has been followed structurally. By examining each stage of the transformation separately, i t has been possible to identify the different structural mechanisms that occur as the process proceeds. The recognition of the important role the noncrystalline polymer plays in the process is a particularly significant result of the study.

INTRODUCTION

Traditionally, elastic polymers have been characterized as rubbers and their mechanical properties have been explained by the entropic mechanisms of rubber elasticity theory. Over the past 15 years, however, growing interest has developed in a new class of crystalline polymer structures that manifest elastic behavior. The elasticity of these hard elastic polymers seems to be primarily controlled by energetic rather than entropic mechanisms. Isotactic polypropylene, poly- ethylene, polybutene-1, poly-4-methyl pentene, acetal copolymer, and polypi- valolactone all produce hard elastic structures.

At least four basic structural models have been proposed to explain the elastic properties of hard elastic polymers. These consist of (a) a leaf-spring involving the elastic bending of lamellae; (b) reversible shear of lamellae between fixed tie points3; (c) a general model based on a change in entropy in the inter- molecular layer and an increase in surface energy during extension4; and (d) a combined structural model that attributes the stress for extension to the pulling of fibrils from lamellae and the retractive force resulting in the particular elastic properties to surface energy and entropy effects in the fibril^.^ These four models are not consistent and predict different mechanical properties and structures in the strained state. Much of this material has been reviewed in the litera- ture.6

The present study examines the formation of a high strength elastic polypropylene fiber. The process is viewed as a series of stages to a final product. The structural changes that occur in the fiber as it passes from one stage to the next is monitored. From this information, a picture emerges of the structural elements that control the elastic behavior, how they rearrange during processing, and how they control the resulting mechanical properties. Emphasis is placed on fiber strength and elastic recovery, as these properties were the primary concern of the study.

* Present address: School of Chemical Engineering, Georgia Institute of Technology, Atlanta, Georgia 30332.

Journal of Polymer Science: Polymer Physics Edition, Vol. 17,535-568 (1979) Q 1979 John Wiley & Sons, Inc. 0098-1273/79/0017-0535$01.00

536 SAMUELS

STRUCTURAL ASPECTS OF HARD ELASTIC POLYPROPYLENE FIBER FORMATION

Both hard elastic polypropylene films and fibers are produced by essentially the same process. That is, the polymer is melt extruded at some melt temperature, drawn down and stress crystallized to a moderate orientation, and then annealed at some temperature for a given time. The morphological structure of these melt drawn and annealed samples results in materials having a high elastic recovery and low strength (1.2-1.5 g per denier). This is illustrated in Figure 1, where the elastic recovery (obtained at both 15% and 50% room-temperature extension) of a series of spun and annealed elastic fibers is plotted. It can be seen from Figure 1 that both annealing temperature and annealing time strongly influence the elastic recovery of the final hard elastic fiber, the best recoveries occurring at the highest annealing times and temperatures shown, with a minimum annealing temperature for good recoveries of 130°C.

The internal structural characteristics of these spun elastic fibers are listed in Table I. The experimental procedures used to obtain these data have been described el~ewhere.~-~ Here p is the fraction of crystals in the fiber, fc and f A M

are the crystalline and noncrystalline orientation functions of the fiber, $A is the fraction of crystal chain axes oriented perpendicular to the fiber axis, f b is the orientation of the b crystal axis with respect to the fiber axis, and AT is the measured birefringence of the fiber.

The orientation function f p is a measure of the average angle the subject axis makes with respect to the fiber axis: fp = (3(cos26) - 1)/2, where 0 is the angle between the subject axis and the fiber axis, and p refers to the phase. Thus, f c is the orientation the molecular chain axis of the crystal ( c axis) makes with respect to the fiber axis, while f A M is the orientation the molecular chain axis in the noncrystalline region makes with respect to the fiber axis. The orientation function f p can vary from a value of +1.0 if all of the chains are oriented parallel to the fiber axis, to a value of zero if the chains are randomly oriented in the fiber, to a value of -0.5 if the chains are oriented perpendicular to the fiber axis.

Figure 2 shows an orientation function triangle diagram for the crystal b and

Extension Annealing temp. 15% 50% ("(3 -- 0 0 150 A A 131 D O 120

0 10 20 30 ANNEALING TIME [MINUTES)

Fig. 1. Effect of annealing conditions on room-temperature elastic recovery of spun elastic fiber.

b l

HIGH STRENGTH POLYPROPYLENE

FleER AXIS t

FIBER AXIS t

a'- A XIS C-AXIS

XIS a'-AXIS fC

I b- h BEFORE ANNEALING AFfER ANNEALING

HARD ELASTIC POLYPROPYLENE

XIS a'-AXIS fC

1.0 b-

BEFORE ANNEALING AFfER ANNEALING

HARD ELASTIC POLYPROPYLENE

+ 'b 1.0

537

Fig. 2. Effect of fiber processing on the crystal orientation of isotactic polypropylene. 0, Annealed elastic fiber; X, melt drawn spun fiber precursor; O , fibers drawn 90°C; A, fibers drawn 90°C and heat set under tension a t 140OC; - - - -, cold drawn films; and V, elastic film.

TABLE I Structural Characteristics of the Hard Elastic Fibers

Sample No. B fc (A V ) f A M $A - f b AT x 10"

0.69 0.71 0.73 0.73 0.70 0.77 0.75 0.72

0.749 0.772 0.773 0.798 0.785 0.763 0.783 0.836

0.331 0.449 0.523 0.476 0.378 0.377 0.500 0.469

0.354 0.348 0.253 0.269 0.296 0.299 0.241 0.222

0.466 0.475 0.444 0.464 0.467 0.456 0.440 0.470

21.2 23.8 24.9 24.6 22.8 22.3 25.9 24.8

c axes in isotactic polypropylene fibers which have undergone different physical treatments. The zero point in the triangle represents random orientation. At each apex, the particular crystal axis will be parallel to the fiber axis, while a point along a side of the triangle would mean that crystal axis is oriented perpendicular to the fiber axis. The dashed line from 0 to C 11 represents the path taken during

538 SAMUELS

cold drawing of isotactic polypropylene film. That is, as the film is drawn, the c axis of the crystal orients parallel to the fiber axis while the a’ and b axes of the crystal vary randomly (line falls equidistant from both sides of the triangle). As can be seen from Figure 2, fibers drawn on a fiber line at 90°C behave in a manner similar to that observed for cold drawn films. If the same fibers are heat set under tension at 140°C then the b axis of the crystal orients perpendicular to the fiber axis faster than the a’ axis, as the c axis goes paralleL7 In fibers spun with a moderate degree of melt orientation the b axis of the crystal is almost perpendicular to the fiber axis (see ref. 8, p. 38), while the a‘ axis is not. This can be seen in Figure 2 where the spun melt drawn precursor fibers all fall on the line very close to the b side of the orientation function triangle. Annealing of these fibers to obtain the hard elastic fiber results in an increase in the c-axis orientation of the fibers while retaining the highly perpendicular b-axis orientation. This suggests the increased c-axis orientation occurs by a rotation of the a’-crystal axis around the b axis to yield more c-axis oriented crystals (see schematic of this process in Fig. 2). An alternative explanation would require melting of the a’-axis oriented crystals with subsequent recrystallization as c-axis oriented crystals. This seems less likely than crystal rotation in light of the morphological results reported below.

The wide-angle x-ray diffraction pattern (WAXS) from both a melt drawn fiber and its elastic counterpart after annealing are shown in Figure 3. A schematic representation of a typical WAXS pattern from a cold drawn fiber is compared with one from an annealed or spun pattern in the same figure. The important feature to note is that the 110 and 130 reflections appear only in the equator of the cold drawn fiber pattern while they also appear along the upper layer line in the spun and annealed patterns. This means there are no a’-axis oriented crystals parallel to the fiber axis in the cold drawn fibers, but there is a fraction of these crystals present in the spun and annealed fibers.

When the a‘ axis of the crystal is oriented in the fiber axis direction then, by necessity, the chain axis (c axis) of the crystal is oriented perpendicular to the fiber axis (see Fig. 2). Since the present models for elastic behavior consider the primary elastic mechanism a consequence of separation of lamellae which have the crystal chain axis oriented in the deformation direction, it is important to know the fraction of crystals present in the polymer which do not satisfy this criterion and how they behave under deformation. The fraction of crystals oriented with the a’ axis parallel to the fiber axis (e.g., c-axis oriented perpendicular) is defined as $A and can be obtained from analysis of the WAXS pattern (ref. 8, p. 144). Figure 4 shows the relation between fc and $A for fibers produced under different processes. Thus the 90°C drawn fibers have about 20% a’-axis oriented crystals in the low draw sample which is quickly pulled out with con- tinued drawing. Heat setting the drawn fiber under tension slows down this process and allows some of the drawn out a’-axis oriented crystals to recover. Melt draw down produces a high fraction of a’-axis oriented crystals (as much as 50% of all the crystals in the sample) which reduces to between 20% and 35% with subsequent annealing to hard elastic fibers. Thus the hard elastic fiber has a large amount of both c-axis oriented crystals and a’-axis oriented crystals.

Not only does the crystal orientation change when the spun fibers are annealed to hard elastic fibers, the noncrystalline region changes as well. In the spun fibers

HIGH STRENGTH POLYPROPYLENE 539

B E F O R E A N N E A L I N G A F T E R A N N E A L I N G

A B

c - A x i s Oriented =-Axis Orientcd

Refleclionr

COLD DRAWN POLYPROPYLENE FIBER ANNEALED OR SPUN POLYPROPYLENE FIBER

Fig. 3. Wide-angle x-ray diffraction of elastic polypropylene fiber. The lower portion shows a schematic representation of polypropylene x-ray patterns.

AM values range from about 0.15 to about 0.30. On annealing, the noncrystalline orientation increases and the AM values for the hard elastic fibers range from 0.33 to 0.52 (see Table I). Thus, annealing of the melt drawn fibers results in a further increase in both the crystal and noncrystal orientation of the already moderately oriented structures with a concurrent decrease in the fraction of a’-axis oriented crystals.

Rotation of pairs of a’-axis oriented crystals around their b axis to form c-axis oriented crystals would lead to perpendicularly oriented noncrystalline chains in the intrapair region. The observed increase in noncrystalline orientation when spun fibers are annealed to form hard elastic fibers suggests plastic deformation of the noncrystalline chains is occurring in the lamella surrounding region. This would impart increased stiffness to the region around the rotating crystals. Since plastic deformation of the less rigid noncrystalline region, and not elasticity, is normally associated with isotactic polypropylene deformation processes, it may be the added stiffness of the surrounding noncrystalline region combined with the lower surface energy required to separate the wide lamellae, which imparts the unique elastic behavior to these fibers.

540 SAMUELS

M I C R OVOID S CON TROLL1 NG

POST DRAWN AND HEAT SET MELT DRAWN

0.6

X

0.4' 1 . 1 I I I I I I

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.0 (PA

Fig. 4. Relation between the crystal orientation function and the fraction of c-axis chains oriented perpendicular to the fiber axis for different fiber processes. 0, annealed elastic fiber; X, melt drawn precursor; 0, fibers drawn 90°C; A, fibers drawn 90°C and heat set under tension at 140°C; and V, elastic film.

STRUCTURAL CHANGES DURING EXTENSION OF HARD ELASTIC POLYPROPYLENE

Once the hard elastic material has been fabricated, then it must undergo some extension to demonstrate its elasticity. Hard elastic isotactic polypropylene fibers have been limited to low strength (<2 g per denier) and high recovery [(go-loo)% after 50% room-temperature extension]. Figure 5 shows the room- temperature stress-strain curve of a hard elastic isotactic polypropylene fiber that has a breaking strength of 1.36 g per denier and an elastic recovery of 96% a t 5w0 room-temperature extension. The sample has two yield points, one in the vicinity of (5-lo)% extension and another in the neighborhood of 6Wo extension. Obviously, several structural processes must be occurring in the hard elastic fiber during room temperature deformation.

It is generally accepted in all of the structural models of hard elastic fiber de- formation14J0 that the separation of crystal lamellae leads to microvoid formation. This is evidenced by the high intensity of diffuse small-angle x-ray scattering (SAXS) observed from the sample during deformation. Figure 6 shows the SAXS patterns from elastic polypropylene film as a function of room temperature extension. Initially, the scattering intensity is low. At 5% extension, the scattering has increased manyfold in the vertical direction, signalling the onset of horizontal microvoid formation. These microvoids have their long axis oriented perpendicular to the deformation direction. As extension proceeds, a horizontal streak begins to appear on the SAXS pattern, becoming quite prominent by (40-60)% extension. This streak signals the formation of vertical microvoids, which have their long axis oriented parallel to the deformation direction. Vertical microvoids are probably generated by the yielding of horizontal microvoids with subsequent extension in the deformation direction.

Microvoids are not the only voids formed during room-temperature extension of hard elastic polypropylene. Figure 7 shows the small-angle light-scattering


1.2 c NONCRYSTALLINE DEFORMATION LONGlTUDl NAL VOIDING SECOND YIELD , Eb = 4.41

db'1.36

0.4

0 -ROTATES TO C-AXIS f', GOES NEGATIVE

1 I I I I I I I I 0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 e

Fig. 5. Stress-strain curve of elastic polypropylene fiber; R = 50%/min.

B

(SALS) patterns obtained from dry elastic polypropylene film at different extensions. The V v SALS patterns are sensitive to density fluctuations in the film, and the changing pattern with extension shows the presence of macrovoids. At 10% extension the VV SALS pattern is extended vertically indicating the macrovoids are oriented with their long axis oriented perpendicular to the draw direction. By (40-60)% extension, a horizontal streak begins to appear indicating the yielding of these macrovoids to form vertically oriented macrovoids (vertical cracks parallel to the deformation direction). If a drop of immersion oil (no = 1.515) is placed on the oriented film, it fills all of the voids with a polymer matching refractive index fluid, and the macrovoid V v SALS pattern disap- pears.

Figure 8 shows light microscope pictures of Pt shadowed carbon replicas of a hard elastic polypropylene fiber both before extension, after 50% extension, and finally after relaxing and reheating. The horizontal macrovoid (macrocracks) are easily seen in the extended sample. Also present are some smaller vertical macrovoids. The fiber before extension and after relaxation and reheating is seen to be free of macrovoids. Thus, upon room temperature extension hard elastic polypropylene forms both macrovoids and microvoids.

Void formation during room-temperature extension is not unique to hard elastic polypropylene, however; Figure 9 shows macrovoid formation in one of the 90°C drawn fiber series samples as seen in the dark field microscope7 during room temperature extension of the fiber. As can be seen from Figure 2, the 90°C fabrication draw temperature corresponds to a cold drawing process for this fiber (points fall along the dashed line of the orientation function triangle diagram). When these 90°C drawn fibers undergo a room-temperature extension, they

542 SAMUELS

U N S T R E T C H E D

10% E X T E N S I O N



5 % E X T E N S I O N

4 0 % E X T E N S I O N

60% E X T E N S I O N 80% E X T E N S I O N 100% E X T E N S I O N

Fig. 6. Effect of room-temperature extension on the small-angle x-ray scattering of elastic polypropylene film. The film axis direction is vertical.

initially form horizontal macrovoids (Fig. 9,5% extension). By 20% extension, the horizontal macrocracks have yielded to form vertical macrocracks. Although these fibers formed macrovoids they formed no microvoids as seen by SAXS until about 20% extension, and even then very little microvoid formation o~curred .~ Heat setting of these fibers under tension changed the internal structure of the fibers, however, so that they now formed some microvoids upon room-temperature extension. This can be seen from the SAXS pattern of the fibers9 shown in Figure 10. Recognizable horizontal microvoid formation starts as early as 2% extension of these fibers with yielding to form some vertical microvoids at around 20% extension. Heat setting also improved the elastic recovery of the fibers (ref. 8, pp. 208 and 209).

The sonic modulus of hard elastic polypropylene film and fiber is shown in Figure 11 as a function of room-temperature extension. Both materials are behaving in a similar fashion. The sonic modulus is seen to continually decrease


20% EXTENSION 40% EXTENSION

543

60% EXTENSION

80% EXTENSION 100% EXTENSION 100% EXTENSION

Vv SALS PATTERNS ZElSS OIL ADDED

Fig. 7. Effect of room-temperature extension on the V v small-angle light scattering pattern from dry elastic polypropylene film. The polarization and machine direction are vertical.

U N S T R E T C H E D 5 0 % E X T E N S I O N R E L A X E D A N D R E H E A T E D

Fig. 8. Effect of treatment on elastic polypropylene fiber surface.

as a function of room-temperature extension until 60% room temperature extension, after which it begins to increase again. This behavior is similar to that observed for both the 90°C drawn, and the heat set under tension fiber~,~,9 and is primarily a consequence of their changing microvoid and macrovoid character.

The uncorrected and void volume corrected birefringence of both dry and immersion oil wetted hard elastic fiberlo and film is shown in Figures 12 and 13 as a function of room-temperature extension. -The initial birefringence of the hard elastic fiber is seen to be considerably higher than that of the film as a consequence of the different processing conditions. The lower birefringence of the film is consistent with its lower crystal orientation (Fig. 2).

As can be seen from the figures, the birefringence of the sample drawn dry is considerably different from that measured with immersion oil present (see Table 11). This is a consequence of void formation which leads to a form birefringence in the sample. Figure 14 shows the form birefringence of hard elastic fiber as

544

096 EXTENSION

SAMUEL S

20% EXTENSlON 3w, ExrENsioN 409 tXlENSlON

Fig. 9. Dark field microscope pictures of void formation in a nonelastic drawn fiber series sample during room-teniperature extension.

0% E X T E N S I O N 2% E X T E N S I O N

4% E X T E N S I O N 15% E X T E N S I O N 30% E X T E N S I O N Fig. 10. Small-angle x-ray scattering from a nonelastic heat set series fiber during room-tem-

perature extension. The fiber axis direction is vertical.

a function of immersion liquid refr'active index (n,) for different room-temperature extensions of the fiber. With no extension, the birefringence remains constant with n,. For extended fibers, the dry negative birefringence value (at n, = 1.0) decreases with increasing immersion liquid refractive index, ultimately becoming positive. It reaches a maximum positive value at a refractive index n, of -1.51 and then decreases again with increasing immersion liquid refractive index. At the point where the birefringence reached a maximum value n, = np (where np is the polymer average refractive index) and the form birefringence, Aform = 0.


Fig. 11. Sonic modulus vs. elongation for elastic polypropylene film and fiber. 0, film; A, fiber.

It is important to know the refractive index of the immersion liquid at which the form birefringence goes to zero. The total birefringence measured dry is given by the expression

AT,dry = AT,polymer -k Aform

AT,^^^^^^^ is required in order to determine the orientation of the noncrystalline region within the deforming system, since (ref. 8, p. 52)

AT,polymer = BAY, + (1 - P)AO,M~AM

Thus, the proper refractive-index immersion medium must be used with hard elastic isotactic polypropylene if the true birefringence of the polymer is to be obtained.

It is also both of interest and important to know the fraction of voids produced in the sample. This is because the measured thickness used to calculate the birefringence must be corrected for the void fraction in order to obtain the true thickness of polymer the light traversed. Figure 15 is a plot of the void fraction &,ojds as a function of room-temperature extension of both the hard elastic film and fiber. The void fraction was calculated from the expression

&ids = 1 - (D,/D d )

where Dd is the measured diameter of the drawn fiber for a given extension ratio A d ; and D, = DoA;”~, where DO is the measured fiber diameter before extension. This expression assumes that the observed deviation of the fiber diameter after extension, from that predicted for an affine deformation, is due solely to void opening. The void fraction obtained for a given room temperature extension of the fiber or film is shown in Figure 15. The void fraction is seen to increase with extension, until almost 30% of the sample is composed of voids at 100% extension. Figures 12 and 13 show the birefringence, both before and after correction for void formation. The corrected birefringence has been used in all subsequent calculations.

Figure 16 is a plot of the crystalline f c and noncrystalline fAM orientation

TA

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ffec

t of R

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-Tem

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Ext

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(AV

) -f

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0.45

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459

0.44

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0.53

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452

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438

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429

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426

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9.56

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560

8.34

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3.79

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*

541

ROOM TEMPERATURE EXTENSION (%)

Fig. 12. Birefringence of hard elastic polypropylene film during room-temperature extension.

functions for hard elastic polypropylene film as a function of room-temperature extension. It should be noted that the noncrystalline orientation function is initially slightly negative (see also Table 11), becoming increasingly more negative with increasing extension until it reaches a maximum negative value within the first 10% extension; accompanied in this region by a slight increase in crystal c-axis orientation. Negative fAM values a t low extension of isotactic polypropylene have been observed before (ref. 8, pp. 61,120, and 121) and have been structurally associated with recoverable crystal alignment in the initial modulus region of the stress-strain curve. As will be shown subsequently, this is consistent with a crystal rotation process (see also Fig. 2) in which a fraction of the a'-axis oriented crystals rotate around the b axis of the crystal to create more c-axis oriented crystals. The limited void formation in this draw region is consistent with a model of crystal rotation as the controlling process in the low draw region.

Between 10% and 40% extension, the crystal fc orientation remains constant (Fig. 16). The noncrystalline orientation also stays essentially constant in this deformation region, but this is due to a reversal of the direction of orientation change with extension, from increasing negative to increasing positive. Ob- viously, a small fraction of the negatively oriented noncrystalline chains are

548 SAMUELS

0

+ 4

- 1 0 - - 1 5 -

-20 - - 2 5 -

0 10 20 30 40 50 60 70 80

ROOM T E M P E R A T U R E EXTENSION(%) Fig. 13. Birefringence of hard elastic polypropylene fiber during room-temperature extension.

becoming positively oriented. This deformation region is the region where void formation predominates (Figs. 6 and 7), and suggests lamella separation processes with its attendant interlamella noncrystalline orientation are controlling. Noether2 has examined SAXS for room-temperature extensions of 30% and greater. Examination of these data (see Fig. 17) shows that the interlamella spacing of the first-order reflection increases directly with extension up to a maximum extension of -75%, after which it deviates from predictions. Also he observed that the second order spacing has an intensity maximum in the (50- 60)% extension range.2 These observations are consistent with the present observations. Thus, the first yield observed in the stress-strain curve (Fig. 5) seems to signal the change from a predominantly crystal-rotation controlling mechanism, to one dominated by lamella separation.

In the region of 40% to 60% extension (Fig. 5) a second yield is observed. This seems to signal the onset of another deformation mechanism, plastic deformation of the noncrystalline region (Fig. 16) with a concurrent yielding of the horizontally oriented voids to form vertical voids (Figs. 6 and 7). The deviation of SAXS spacings from predictions (Fig. 17) with concurrent decreases in SAXS intensity


r

+ 2 5 t 549

9

"S

Fig. 14. Form birefringence of elastic polypropylene fiber. Extension (%) 0 0 0 6 0 A 10 V 80 + 20 x 100 0 4 0 A 150

of the second order reflection beyond 60% extension supports this conclusion. Figure 18 shows the contribution of each of these processes to the measured birefringence of hard elastic polypropylene film during room-temperature extension.

The significance of the presence of a large amount of a'-axis orientation in isotactic polypropylene has been largely neglected in previous studies of hard elastic polypropylene. This is not surprising since the controlling elastic mechanism in the most important deformation region (>lo% extension) seems to be the lamella separation process. Nevertheless, the presence of a crystal rotation process should not be ignored. Figures 19, 20, and 21 illustrate the nature of this crystal rotation process during room-temperature extension in greater detail. Here the (110) reflection of the WAXS pattern has been de-

550

-0,2!A

-0.3

-0 .4

- 0 . 5

0.3

~ 0 . 2 0

0 > *

8

0.1

N O N C R Y S T A L L I N E 'LA O R I E N T A T I O N -

- I I I I I I I I I I

SAMUELS

0

I I 1 I I I I I I I 0 1 0 20 30 40 5 0 6 0 70 80 90 100

ROOM T E M P E R A T U R E EXTENTION(%)

Fig. 15. Calculated fraction of voids (f#Jv&&) as a function of room-temperature extension of isotactic polypropylene hard elastic fiber and film. 0, Elastic fiber; X, elastic film.

O R I E N T A T I O N +0.4b

- +0.2 +0.3 t

R O O M T E M P E R A T U R E EXTENSION(%) Fig. 16. Crystalline and noncrystalline orientation of hard elastic polypropylene as a function

of room-temperature extension. 0, f c ; A, f A M .

composed (ref. 8, p. 144) in order to examine the orientation function of that fraction, @A, of crystal chain axes that is oriented perpendicular to the deformation direction fc,l (a'-axis orientation), and that fraction & of crystal chain axes that are oriented parallel to the deformation direction, f,, 11 ( c -axis orientation). It is the vectorial sum of both these contributions which produces the total average crystal c-axis orientation, f c (AV) , normally obtained from the diffrac- tometer (Figs. 10 and 19).


R O O M T E M P E R A T U R E E X T E N S I O N ( % ) Fig. 17. SAXS spacings vs. extension of hard elastic polypropylene (ref. 2). 0, Polypropylene

= L = 200 fiber; -, structure model; - - - -, best line through data (Samuels); a = 150 A, I = 50 A; A.

as the vector sum of the individual crystal contributions Figure 19 shows fc ( A V ) , f ( : , I, fc , 11, f b , and fc ( A W c a l c . f c ( A W c a l c is calculated

fc ( A W c a l c = [ ( 6 A f c , I )2 -I- ( 6 c f c , I1 )21 1’2

As can be seen from Figure 19, the crystals that are oriented with their chain axes parallel to the deformation direction fc,ll are highly oriented initially and then decreased slightly in orientation with increasing room temperature extension. The crystals oriented with their chain axis perpendicular to the deformation direction fc , I are highly oriented in the perpendicular direction and become somewhat more oriented in the stretch direction (less perpendicularly oriented) a t extensions greater than 60%. The fc , I values also correspond closely to the f b (b-axis orientation) values indicating these axes vary randomly with a’-axis orientation. Thus, there are two families of highly oriented crystals present in the sample, the orientation of each of which varies very little with room-temperature extension. The observed increasing f c (A V ) value with initial extension of the sample is thus primarily a consequence of the changing relative number of crystals oriented parallel to the deformation direction as compared to those oriented perpendicular.

This effect is illustrated in Figures 20 and 21. In Figure 20, the fraction of crystals oriented with their c axis perpendicular to the deformation direction 4 ~ , and the fraction of crystals oriented with their c axis parallel to the deformation direction 4c = (1 - @ A ) , are plotted against the room temperature ex-

552 SAMUEL S

(CRYSTAL CONTRIBUTION) L X ” X

/- 5A0cfc

(NONCRYSTALLINE CONTRIBUTION)

-5

AT ( D R Y ) c o R R E CT E D

0-0

AFORM ( F O R M C O N T R I B U T I O N )

“0 - 1 0 - 7

x U

- 1 5 -

0-0-

0 10 20 30 40 50 60 70 80 90 100

-20 - -25 -

-30 -

-35

- 4 0 I I

ROOM TEMPERATURE EXTENSION(%) Fig. 18. Birefringence contributions in hard elastic polypropylene during room-temperature ex-

tension.

tension of the film. In Figure 21, the percentage change in the fraction of a’-axis oriented crystals from those originally present before stretching ( A ~ A ) % , is plotted (see also Table 111). Here

( A @ A ) % = [ (dA,O - d’A,X)/4A,O] x 100

where ~ A , O is the initial fraction of a’-axis oriented crystals before stretching and ~ A , X is the fraction of a’-axis oriented crystals present at room temperature extension A.

Before extension, there are approximately equal amounts of a’-axis (fc, I ) and c-axis ( f c , 1 l ) oriented crystals in the film. On initial extension, the amount of a’-axis oriented crystals quickly decreases (Fig. 20), so that, by 10% extension of the film, 20% of the original a’-axis oriented crystals have converted to c-axis oriented (fc,1l ) crystals (Fig. 21). From 10% to 40% room-temperature extension, the fraction of chains oriented parallel and perpendicular to the deformation direction remains essentially constant (Figs. 20 and 21), only 4% of the original, perpendicularly oriented crystals converting to parallel orientation. Beyond


C - A X I S P A R A L L E L O R I E N T A T I O N

” *-

C - A X I S P E R P E N D I C U L A R O R I E N T A T I O N AND f b

-0. $--T-n SI I - Jy----F---K 0 10 20 30 40 50 6 0 70 80 90 100

ROOM T E M P E R A T U R E E X T E N T I O N ( % )

Fig. 19. Orientation function component behavior as a function of room-temperature film extension. 0, f c , l l (c-axis parallel component); 0, f c , l (c-axis perpendicular component); X, fb,’ A,

fc(AV),,i, = [ ( @ ~ f c . i ) ~ + (@cfc,11)211’2; +, fc(AV),,pt.

TABLE I11 Effect of Room-Temperature Extension on the a’-Axis Orientation of Hard Elastic

Polypropylene Film

f c ( A V ) (%) c a 1 c u 1 ate d

Exten- (vector sion f c ( A V ) @A - f C J &A @c f c , II sum)

0 0.453 0.479 0.430 0.000 0.521 0.923 0.523 2 0.440 0.474 0.447 0.010 0.526 0.906 0.522 5 0.500 0.424 0.443 0.115 0.576 0.911 0.557

10 0.536 0.383 0.440 0.200 0.617 0.911 0.587 20 0.549 0.369 0.433 0.230 0.631 0.906 0.594 40 0.550 0.363 0.447 0.242 0.637 0.888 0.588 60 0.577 0.325 0.451 0.322 0.675 0.881 0.612 80 0.542 0.342 0.415 0.286 0.658 0.858 0.582

100 0.564 0.297 0.416 0.418 0.703 0.844 0.606

4090 extension, the fraction converting to parallel orientation increases again so that by 100% extension another 20% of the original a’-axis oriented crystals are converted to c-axis orientation (Table 111).

Examination of the structural data presented in this section suggests there are two families of crystals produced in the initial hard elastic polypropylene. One of these families of crystals (consisting of many of the original c-axis oriented crystals) contributes to the elastic behavior by a microvoid (crystal lamella surface separation) process, while the other undergoes the more usual crystal rotation and plastic deformation processes. On initial room-temperature ex-

554 SAMUELS

0.

0.

0.

,$

0.

0.

I] - @c ( C - A X I S P A R E L L E L )

DEFORMATION D I R E C T I O N

c - A X I S

R O T A T I O N c - A X I S c 2. Z O X O F P E R P E N D I C U L A R CRYSTALS

CONVERTED TO PARELLEL

O R I E N T E D PARELLEL TO THE DEFORMATION D I R E C T I O N

--0 ~ ~ ~ ~ ~ ~ ~ ~ ~ ~ ~ A ~ I ~ ~ c ~ ~ ~ N ------- 0

0

0- '$'A ( c -AX I S P E RP END I CU L A R ) BEFORE E X T E N S I O N AFTER E X T E N S I O N

ROOM TEMPERATURE E X T E N S I O N ( % )

Fig. 20. Effect of room-temperature extension on the fraction of a'-axis oriented crystals in elastic polypropylene film.

50

4 0

30 4 8 a -

20

10

0

( C R Y S T A L R O T A T I O N ) - fAM I N C R E A S I N G N E G A T I V E

F I R S T Y I E L D I N S T R E S S - S T R A I N

P O L Y M E R Y I E L D I N G 0-O - aye----+ ( P L A S T I C D E F O R M A T I O N )

V O I D O P E N I N G f A M I N C R E A S I N G P O S I T I V E

( L A M E L L A S E P A R A T I O N )

fAM C O N S T A N T

I 1. I I

ROOM T E M P E R A T I J R E E X T E N S I O N ( % )

20 40 60 8 0 100

Fig. 21. Change in the percent of a'-axis oriented crystals with room-temperature extension of hard elastic polypropylene film.

tension, a large fraction of the a'-axis oriented crystals rotate around the b axis so that their chain axis is now parallel to the deformation direction. This process is similar to orientation behavior previously observed during high-temperature


extension of films (ref. 8, p. 120). The nondisruptive rotation of the crystals also leads to an increasing perpendicular orientation of the interlamellae noncrystalline chains (Fig. 16). However, as extension proceeds toward ~ W O , the stress increases, and increasing amounts of microvoids begin to compete with the crystal rotation process. The first yield region of the stress-strain curve (Fig. 5) mirrors this change from a crystal rotation process to one where crystallite separation with microvoid formation predominates.

That the region between 10% and (40-60)% room-temperature extension is dominated by void opening processes is reflected in the essentially constant +A,

fc , and f A M values observed (Figs. 16,19,20, and 21) with concurrent increasing contributions from such void sensitive measurements as Aform, SAXS, SALS, and void volume (Figs. 6,7, and 12-15). Beyond (40-60)% extension of these samples, the continually increasing stresses leads to increasing competition from plastic deformation processes and this results in a second yield point in the stress-strain curve (Fig. 5). The increasing contribution of plastic deformation processes is manifested by the increasing orientation of the noncrystalline chains (Fig. 10); the significant increase in the fraction 4~ of a’-axis oriented crystals drawn out to c-axis orientation (Fig. 21); the fact that these crystals are not as well oriented as the earlier rotated crystals (this is seen in Fig. 19, where fb, fc,l,

and f c , I I are all becoming less perfectly oriented); and the yielding of both micro- and macrovoids to form vertically oriented voids (Figs. 6 and 7). These processes are summarized schematically in Figure 22. Of course all three processes, crystal-rotation, void formation by lamella separation, and plastic deformation, are occurring at extensions greater than 6Wo; but at these extensions the plastic deformation processes become increasingly dominant.

E X T E N S I O N E X T E N S I O N D I R E C T I O N D I R E C T I O N

C-AXIS

E X T E N S I O N D I R E C T I O N

E Fig. 22. Schematic representation of the deformation controlling processes during room-tem-

perature extension of hard elastic polypropylene.

556 SAMUELS

HIGH STRENGTH ELASTIC POLYPROPYLENE

In a normal fiber process, the fiber is first melt spun at a certain melt temperature, drawn down, and then quenched. This spun fiber will have a low tenacity and a high extensibility. In order to produce strong fibers, the spun fiber is then drawn at some lower draw temperature to produce an oriented fiber with the desired high strength properties. This was the method used to prepare the 90°C drawn and the heat set fibers, discussed in the section on structural aspects of hard elastic polypropylene fiber formation (Ref. 8, p. 140). Thus, the low strength spun fiber is used as a precursor for production of the high strength drawn fibers.

By analogy then, the low strength elastic fiber should be used as the precursor for production of a high strength elastic fiber since the low strength elastic fiber is simply an annealed spun fiber. The question is how to orient this elastic precursor without destroying its elastic character.

One possible solution is suggested by examining the deformation mechanisms the hard elastic fiber (designated elastic precursor) undergoes during room temperature extension (see the section on structural changes during extension of hard elastic polypropylene). When the precursor elastic fiber is extended beyond 10% extension microvoids open up. It is the reclosing of these microvoids that produces the retractive force that is seen as the fibers recovery. Why not use the fibers own internal retractive force to obtain the desired orientation?

Figure 23 schematically shows how this can be done (see also Fig. 22). The precursor elastic fiber initially contains a large amount of c-axis oriented crystals which on room-temperature extension will separate and produce microvoids. Drawing the precursor elastic fiber to relatively high extensions (100%-200%) will cause extensive microvoid production as well as some plastic deformation. If the fiber is unclamped at this point, strain release by both the large number of highly strained noncrystalline chains between the separated c -axis oriented lamella, and by the strained lamella themselves, will cause the fiber to retract.

S T R A I N E D S T R A I N E D N O N C R Y S T A L L I N E L A M E L L A C H A I N S

C R Y S T A L L A M E L L A c - A X I S O R I E N T E D

I

VOID

NONCRYSTALLINE C H A I N S

C R Y S T A L L A M E L L A R E L A X E D

N O N C R Y S T A L L I N E C H A I N S O R I E N T E D

N O N C R Y S T A L L I N E C H A I N S R E L A X E D

U N S T R E T C H E D ROOM T E M P E R A T U R E H I G H T E M P E R A T U R E S T R E T C H ANNEAL UNDER T E N S I O N

Fig. 23. Schematic representation of the formation of high strength elastic fiber structure.


* v) < 6 o -I w

50

4 0

Thus, little improvement will be imparted to the sample. If the extended fiber is prevented from retracting and the temperature is raised, a different process of recovery would be predicted. Since the melting temperature of a crystal system containing highly strained noncrystalline chains is higher than that containing less strained noncrystalline chains,ll an increase in the temperature of the restrained drawn elastic precursor will soften the less oriented noncrystalline material surrounding the open microvoid. The high retractive force of both the oriented chains within the microvoid, and the strained lamella, will no longer be prevented by a rigid surrounding matrix from closing the microvoid. Instead, they will close the microvoid by drawing out the now softer surrounding material. Thus, the closing of the microvoids within the restrained fiber will act in a similar manner to the surrounding matrix as a drawing process does on a nonelastic fiber line. Orientation of the polymer in the elastic system however, will occur by drawing between closing microvoids (Fig. 23). For example, if a precursor elastic fiber is used which when drawn to 100% extension at room temperature, recovers 96% of that extension by microvoid closing when released, it will effectively draw the softer polymer that surrounds the microvoid the same 96% extension during void closing when held at fixed length at an elevated temperature. Thus, heat setting the elastic precursor fiber at a fixed extension should lead to a high strength elastic fiber. It is interesting to note that one of the fabrication procedures used by Noether and Brody12 to produce “low modulus” hard elastic polypropylene fibers involved just such an annealing of drawn precursor elastic fiber under tension. Examination of their data shows this procedure produced high strength elastic fiber as predicted above.

The precursor elastic fiber chosen for the present study had 90% recovery at 50% room-temperature extension (Fig. 24). The fiber was extended between

- 0.077 \ \*\ \ ‘*\ ORAWN AND HEAT S E T

0.018 0.000 \*\ NORMAL FIBER (2.9g/DEN.)

Q A = 0 . 0 9 4 4‘

- A, ’.

90°C DRAWN NORMAL FIBER = 0 . 0 4 2 (3.3g/OEN.) -

I I I I I I I I

ROOM TEMPERATURE CYCLED EXTENSION(%)

Fig. 24. Elastic recovery under cycled extension of several types of isotactic polypropylene. - X, Elastic film; - 0, elastic precursor fiber; - 0.100% extension 30 min: - A, 150% extension 30 min; - - - - +, 200% extension 30 min; - - - - 0, drawn and heat set fiber 2.9 (g/denier); and - - - - A, 90°C drawn fiber (3.3 g/denier).

TA

BL

E I

V

Stru

ctur

e and

Pro

pert

ies o

f Hig

h St

reng

th E

last

ic F

iber

s (H

eat S

et a

t 16

OOC

for

30 m

in)

Roo

m-

Sam

ple

tem

pera

ture

T

enac

ity

No.

ex

tens

ion

P fc

-f

b

AT

X lo

3 f

AM

@

A

-f

Gl

fc, /I

x

b

(g/d

enie

r)

L (A

) 1

0 0.

717

0.77

9 0.

447

18.6

2 0.

139

0.23

9 0.

302

0.89

6 5.

41

1.36

17

0.0

2 10

0 0.

710

0.87

4 0.

453

24.1

5 0.

350

0.07

7 0.

362

0.91

1 2.

00

2.96

18

5.8

3 15

0 0.

730

0.93

2 0.

469

26.6

2 0.

420

0.01

8 0.

500

0.93

6 1.

88

3.93

18

9.0

4 20

0 0.

707

0.88

9 0.

453

29.6

3 0.

645

0.00

...

0.92

5 1.

49

5.55

18

6.6


20 I 1

1 0 - z 6 - 2 5 -

e 4 -

9 - 8 - - 7 -

w

m - 3 - > + - 2 2 - r w +

1, . 9 - .a - .7 - .6 - - 5 t-

0.11 I 1 1 1 1 1 1 1 1 1 1 - 0 . 2 - 0 . 1 0 0.1 0 . 2 0.3 0 . 4 0.5 0.6 0.7 0.8 0 . 9 1 . 0

Fig. 25. Relation between the tenacity and the orientation functions for high strength elastic polypropylene (the solid fa^ line is that predicted from 30 different nonelastic samples of polypropylene).

clamps either 100%, 150%, or 200% at room temperature and then placed in an oven at 160°C for 0.5 hr. The long time was chosen in order to ensure closure of both the microvoids and macrovoids in the samples. The mechanical properties and internal structure of these high strength elastic (HSE) fibers were then determined (see Table IV).

The tenacity of drawn isotactic polypropylene fibers and films is determined by the orientation of the noncrystalline chains in the fiber (ref. 8, p. 198). It is of interest to see if the same relationship between tenacity and noncrystalline orientation exists between these oven set HSE fibers and normally drawn isotactic polypropylene fibers and films. Figure 25 shows that the same relationship does exist between these oven set HSE fibers and the normal drawn polypropylene fibers and films. It further shows that the tenacities are not a function of their crystal orientation. It should be noted that oven setting of the strained elastic precursor achieved tenacities as high as 5.5 g per denier.

The stress-strain curves for the elastic fibers are shown in Figure 26. The (40-60)% extension yield region of the elastic precursor (see also Fig. 5) is still present as a rate of change in slope of the stress-strain curve in the HSE fibers. A small yield is also present at 5% extension of the 3-g/denier HSE fiber but this yield point is missing from the 4- and 5.5-g/denier HSE fibers. The unusual nature of the HSE fiber yield behavior is more easily recognized when their stress-strain curves are compared with drawn (90OC) and drawn and heat set (14OOC) fibers of essentially the same tenacity (Figs. 27-29). The nonelastic fibers have a sharp yield in the vicinity of (20-30)% extension, while the HSE fibers show simply a continuous change in slope in the (40-60)% extension region. The (20-30)% extension region of the nonelastic fibers is where plastic deformation completely dominates their deformation b e h a ~ i o r . ~ , ~ In the HSE fibers, the (40-60)% region is the region in which plastic deformation processes are making a greater and greater contribution toward the extension behavior with

560 SAMUELS

6 . 0

0- Eb = 4 . 4 1 / ~ - - - o

0 I I I I 0 0 . 5 1 . o 1 . 5 2.0

Fig. 26. Comparison of the stress-strain curves from different strength elastic polypropylene fibers. 0, Original elastic fiber; 0 , 3 0 min at 100% extension; ., 30 min at 150% extension; and A, 30 min at 200% extension.

6 . 0

5 . 0

- r

4 . 0 - 1

D

-

-

-

Fig. 27. Comparison of the stress-strain curves from drawn and heat set fibers and elastic fibers having tenacities of about 3.0 g/denier (R = 50%/min). A, Drawn fiber (3.3 g/denier); 0, heat set fiber (2.9 g/denier); 0, original elastic fiber; and O,30 min at 16OOC (2.95 g/denier).

increasing extension and thus the transition is a gradual, rather than a sharp one.

The yield observed at (&lo)% extension of the elastic precursor was identified with initial recoverable alignment of the a’-axis oriented crystals (see the section on structural changes during extension of hard elastic polypropylene). Heat setting of the strained elastic precursor rapidly orients the a‘-axis oriented crystals so that the sample annealed at 100% extension (sample 2, Table IV) has only 8% of the crystals a’-axis oriented, while samples 3 and 4 are essentially free of a’-axis oriented crystals. The suggestion of only a small yield in sample 2 and no yield in samples 3 and 4 at (&lo)% extension is consistent with these results.


5

Fig. 28. Comparison of the stress-strain curves from drawn and heat set fibers and an elastic fiber having tenacities of about 4.0 g/denier (R = 50%/min). A, drawn fiber [4.1 (g/denier)]; 0 , heat set fiber [3.7 g/denier)]; 0, original elastic fiber 11.3 (g/denier)]; and 0 , 1 5 min at 160°C [4.2 (g/denier)].

Fig. 29. Comparison of the stress-strain curves from drawn and heat set fibers and elastic fiber having tenacities of about 5.5 g/denier ( R = 50%/min). A, Drawn fiber (5.7 g/denier); 0, heat set fiber (5.2 g/denier); 0, original elastic fiber; 0,30 min at 160°C (5.6 g/denier); and B, 60 min at 160°C (5.8 g/denier).

The results are also consistent with the observed elastic recovery of these fibers. The elastic recovery of the fibers is defined as the fractional length the fiber re- covered from its initial extension. This was obtained by cycling a given fiber in the Instron tensile tester through successively increasing cycles to 50% extension. The extensional strain of a given cycle is given as

EE = (LE - Lo)/& where LE is the length of the fiber at its fullest extension during the cycle and LO is'the initial length of the fiber before any cycle extension. The final strain after a given cycle is given by

EF = (LF - Lo)/Lo

562 SAMUELS

where LF is the length of the fiber at the end of a complete cycle, and Lo is the same as above. The percent elastic recovery RE(%) of a given cycle, is then given by the expression

R E ( % ) = [(EE - CF)/EE] x 100

The elastic recovery RE(%) of the elastic precursor, the HSE fibers, a 2.9-g/ denier drawn nonelastic fiber, and a 3.3-g/denier heat set nonelastic fiber, are plotted against the percent room-temperature cycled extension in Figure 24. The differences between the elastic recovery of these fibers in the low cyclic extension region (5-10%) is in general agreement with the previous conclusion (see the section on structural changes during extension of hard-elastic polypropylene) that the fraction of a’-axis oriented crystals present in the sample has a significant influence on the observed recovery in this deformation region. The oven heat setting of the restrained elastic precursor is seen to have led to a less elastic, though stronger, HSE fiber. These HSE fibers are seen, in turn, to have better elastic recovery, at cyclic extensions greater than 20%, then nonelastic fibers of about 3 g per denier strength.

Since the noncrystalline orientation seems to play an important role in the deformation of elastic polypropylene a t room-temperature extensions greater than 20%, the elastic recovery of the nonelastic and elastic fibers is replotted in Figures 30-32 as a function of the measured noncrystalline orientation in the fibers before any cycling for each cycle extension. For any given fiber, the elastic recovery is seen to decrease with cycle extension. For the drawn (90OC) nonelastic fibers (Fig. 30), the elastic recovery a t a given extension increases with increasing noncrystalline orientation of the fibers. For the heat-set nonelastic fibers (Fig. 31), the elastic recovery a t a given extension is more complex. It is sensitive to the noncrystalline orientation in the low orientation range and almost

01 I I I I I I I I 0 0.1 0 . 2 0.3 0.4 0.5 0.6 0.7 0.8 C

AM

9

Fig. 30. Elastic recovery of drawn (90°C) isotactic polypropylene fibers under cycled extension as a function of their noncrystalline orientation.


100

90

- w. 80 - 2. = > 0 v w

w 70

LL 60 " c m -

50 W

4 0

I

- x I

C Y C L E E X T E N S I O

SPUN F f B E R P R E C U R S E R 0 2 0 4 - 0 - 5 +

- /;$-

l o 1 5 2 0

2 5

30

35

4 0 4 5 50

30 t 0 1 I I I I I I I I

0 0 . 1 0 .2 0.3 0.4 0.5 0.6 0.7 0 . 8 0

f~~

Fig. 31. Elastic recovery of heat set fibers under cycled extension as a function of their noncrystalline orientation.

insensitive to the noncrystalline orientation a t orientations greater than f A M = 0.4. There is some increase in RE with f A M in the (15-20)% cyclic extension range at f A M > 0.4 but this reverses by 30% cyclic extension, so that the elastic recovery actually decreases with increasing noncrystalline orientation at extensions greater than 30% cyclic extension for the heat set nonelastic fibers with f A M > 0.4.

The elastic recovery of the hard elastic fibers (Fig. 32) seems very different from the nonelastic fibers a t first glance, however, their behavior actually has many similarities to the heat set nonelastic fiber behavior. Thus, at f A M 1 0.35 the elastic reocvery is insensitive to f A M in the cyclic extension region of (5-35)%. Also for cyclic extensions greater than 35%, the elastic recovery decreases with increasing noncrystalline orientation just as with the nonelastic heat set fibers. The HSE fibers have a smaller elastic recovery range over which they change than does the nonelastic fibers so that there is a significant difference between the two types of fibers, but the general character of their behavior in the region f A M > 0.35 is similar. The most striking difference between the elastic and nonelastic recovery curves is in the behavior of the precursor. Thus, the spun fiber from which the drawn (goo) and the heat set nonelastic fibers were obtained has such poor elastic recovery that almost any orientation seems to improve its behavior. On the contrary, the hard elastic precursor has excellent elastic properties and orienting this material diminishes its behavior.

Figure 33 shows the elastic recovery at 50% cyclic extension of the nonelastic and elastic fibers as a function of their noncrystalline orientation. It also shows that the elastic recovery a t 50% extension, like the tenacity, is independent of the crystal orientation in the HSE fibers. This figure illustrates two important differences between the elastic and nonelastic fibers: (a) the elastic recovery of the elastic fibers is superior to that of the nonelastic at all orientations, and

564 SAMUELS

100

90

80

bp

* Y > 0 u Y

- = 70

p: 60 V

C Y C L E EXTENSION

%

u

* VI

U _1 Y

-

Fig. 32. Elastic recovery of elastic polypropylene fibers under cycled extension as a function of their noncrystalline orientation.

(b) while orientation of the nonelastic fiber leads to improvement in the elastic recovery at 50% extension over that of the spun precursor; increased orientation of the HSE fiber leads to reduced 50% cyclic elastic recovery over that of the precursor elastic fiber. It should be noted (Figs. 32 and 33) that the elastic recovery of the HSE fibers at 50% cyclic extension is a linearly decreasing function of the noncrystalline orientation in the fiber. Thus, the elastic recovery at 50% cyclic extension decreases as the tenacity increases (Fig. 25) in these HSE fibers due to the influence of the noncrystalline orientation in each. A plot of tenacity against the elastic recovery at 50% cyclic extension is plotted in Figure 34. This further illustrates the intimate relationship between these two physical measurements.

In general then, the observed properties of the HSE fibers are consistent with the structural models presented in Figures 22 and 23. That is, three regions of deformation behavior appear in the HSE fibers just as in the hard elastic precursor. In the initial region [(0-lo)%] extension, a tensile yield can occur in the HSE fibers if the fraction of a'-axis oriented crystals .is high enough. The room temperature elastic recovery of isotactic polypropylene fibers generally seems also to be controlled by a'-axis orientation in this low extension region as evidenced by the cyclic recovery of the elastic and nonelastic fibers shown in Figure 24. The elastic recovery of the HSE fibers is essentially independent of their initial noncrystalline orientation in this region (Fig. 32).

The behavior of the HSE fibers in the region of (10-35)% room-temperature extension is also consistent with the structural models presented in Figures 22 and 23. Thus, just as with the hard elastic precursor, the yield on the stress- strain curve does not occur until after this extension (Fig. 26), and appears as


100-

90

80 w 0 Ln

Y

" 70 - 2 60 W

> 0 V

p:

V

w 50

,-. k 4 0 - 4 -1 W

t p 30

20

10

- ELASTIC P R E C U R S E R -

0

-

-

- NON-ELASTIC FIBERS -

- f NON-ELASTIC DRAWN(90"C) FIBERS SPUN F I B E R

-

I I I I I I I I I I I I -0.2-0.1 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0

as a function of their noncrystalline orientation.

c

6.0- - z W n

p: 5 . 0 - W n.

4.0- 4

u7 a

* 3.0- - + V 4

I

2 . 0 -

1 . o

+

ELASTIC FIBERS

NONELASTIC O R A W N (900c) F I B E R S i\NoNE;;;\; HEAT SET (140°C)

I I I I I I I I it'

a smooth transition over a wide deformation range [(35-60)% extension] as compared to the sharp yield transitions of comparable strength nonelastic fibers, demonstrating the influence of microvoid formation in this [( 10-35)%] extension region (Figs. 27-29). Similarly, the elastic recovery of the HSE fibers is independent of the cyclic extension in the (10-20)% extension region and then decreases slowly beyond this, demonstrating the strong influence of microvoid

566 SAMUELS

AM Fig. 35. Relation between the yield strength nY, and the noncrystalline orientation f,q.g for high

strength elastic polypropylene.

recovery processes in this region and the increasing influence of plastic deformation. Again, in this [(lo-35)%] cyclic extension region, the recovery is independent of the initial noncrystalline orientation in the fiber.

Finally, the behavior of the HSE fibers in the deformation region beyond 35% extension is also consistent with the structural models of the behavior of the hard elastic precursor shown in Figures 22 and 23. Thus, plastic deformation processes and the initial noncrystalline orientation of the HSE fibers plays an increasing role. For example, the yield region of the stress-strain curve appears over this extension region [ (35-60)%] indicating a transition in deformation processes. The yield strength of the elastic precursor and HSE fibers is a function of the initial noncrystalline orientation in the fibers (Fig. 35), as is also the tenacity (Fig. 25). Finally, the cyclic elastic recovery is an increasingly de- pendent function of the initial noncrystalline orientation in the HSE fibers (Fig. 32), and in fact, is a linear function of fAM at 50% cyclic extension of the fibers (Fig. 33).

CONCLUSIONS

The thrust of the present study has been to follow structurally the transformation of a spun fiber into a high strength elastic fiber. By studying each step in the transformation of a spun fiber into a high strength elastic fiber separately, it has been possible to isolate the different structural transformations that occur in the fiber as it progresses through the process. The major conclusions of the study are the following:

Hard Elastic Polypropylene Formation

(a) The minimum annealing temperature required to transform a moderately oriented spun polypropylene fiber into a hard elastic fiber is 130°C.


(b)

(c)

The highest annealing temperature used in the study (150°C) yields hard elastic fibers with the best elastic recovery.

The structural transformation from a spun fiber to a hard elastic fiber involves a decrease in the fraction of a’-axis oriented crystals, an increase in the chain axis orientation of the crystals, and an increase in the orientation of the noncrystalline chains. These observations suggest the elastic character of the hard elastic fibers may be a consequence not only of the size and orientation of the crystal lamellae, but of the stiffness of the noncrystalline matrix as well.

Hard Elastic Polypropylene Room-Temperature Deformation

(a) There are three room-temperature deformation regions in hard elastic polypropylene. Each region is dominated by a different structural mechanism.

(b) The initial deformation region [(0-lo)% extension] is dominated by crystal rotation processes; in particular, the rotation of the a’-axis oriented crystals around their b axis, to form c-axis oriented crystals.

The intermediate deformation region [( 10-40)% extension] is dominated by void formation. This is the region where lamella separation processes are controlling.

The final deformation region (>40% extension) is the region in which plastic deformation processes become increasingly dominant.

(c)

(d)

Formation of High Strength Elastic Fibers

(a)

(b)

Heat-setting of drawn and restrained hard elastic fibers (designated precursor elastic fibers) in an oven at 160°C for 30 min yields HSE fibers.

Tenacities as high as 5.5 g per denier were achieved and the noncrystalline orientation in the HSE correlates with tenacity according to the same form as found for nonelastic polypropylene fibers and films.

(c) (d)

(e)

The tenacity is independent of the crystal orientation. The HSE fibers do not show a sharp yield in their stress-strain curve but

instead a slow change in slope with extension is observed. The HSE fibers have significantly better elastic recoveries from room

temperature extensions greater than 20%, than observed for nonelastic fibers of the same tenacity.

The elastic recovery at 50% room-temperature extension of the HSE fibers is directly proportional to the initial noncrystalline orientation in the fiber and independent of the crystal orientation in the fiber.

(g) Three regions of deformation behavior appear in the HSE fibers which correspond mechanistically to similar regions observed in the precursor elastic fiber.

A particularly important conclusion of this study has been the identification of the role the noncrystalline region plays in elastic fiber properties. The noncrystalline region has been largely neglected by other investigators as a significant contributor to elastic fiber properties. This study shows that such neglect was unfortunate and that an understanding of the role of the noncrystalline region is essential if elastic fiber behavior is to be fully understood.

(f)

568 SAMUELS

The author wishes to acknowledge the assistance of Dr. J. P. Modrak for supplying the elastic precursor fibers, R. H. Kridler for the birefringence and x-ray data on the HSE fibers, and C. T. Hightman for the stress-strain and recovery data on the HSE fibers.

References

1. B. S. Sprague, J. Macromol. Sci., Phys., 8 (1,2), 157 (1973). 2. H. D. Noether and W. Whitney, Kolloid Z. Z. Polym., 251,991 (1973). 3. E. S. Clark, Structure and Properties of Polymer Films, R. W. Lenz and R. S. Stein, Eds.,

4. D. Goritz and F. H. Muller, Colloid Polym. Sci., 252,862 (1974). 5. M. Miles, J. Petermann, and H. Gleiter, J. Macromol. Sci., Phys., 12 (4), 523 (1976). 6. S. L. Cannon, G. B. McKenna, and W. 0. Statton, J. Polym. Sci., Macromol. Reu., 11,209

7. R. J. Samuels, J. Polym. Sci., A-2, 6,2021 (1968). 8. R. J. Samuels, Structured Polymer Properties, Wiley, New York, 1974. 9. R. J. Samuels, J. Polym. Sci., Part C, 20,253 (1967).

Plenum, New York, 1974.

(1976).

10. I. K. Park and H. D. Noether, Colloid Polym. Sci., 253,824 (1975). 11. R. J. Samuels, J. Polym. Sci., Polym. Phys. Ed., 13,1417 (1975). 12. H. D. Noether and H. Brody, Text. Res. J., 46,467 (1976).

Received February 17,1978 Revised July 12,1978

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High strength elastic polypropylene

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