Oriented structures based on flexible polymers : drawing behaviour and properties

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Oriented structures based on flexible polymers : drawing

behaviour and properties

Citation for published version (APA):

Bastiaansen, C. W. M. (1991). Oriented structures based on flexible polymers : drawing behaviour and properties. Technische Universiteit Eindhoven. https://doi.org/10.6100/IR356115

DOI:

10.6100/IR356115

Document status and date: Published: 01/01/1991

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Oriented Structures Based on Flexible polymers

Drawing Behaviour and Properties

7.99

3.28

1.13

0 0 0

3.28

9.92

2.14

0 0 0

c

pq

=

1.13

2.14

316

0 0 0

GPa

0 0 0

3.19

0 0 0 0 0 0

1.62

0 0 0 0 0 0

3.62 Cees Bastiaansen

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Oriented Structures Based on F'lexible Polymers

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Oriented Structures Based on Flexible Polymers

Drawing Behaviour and Properties

Proefschrift

ter verkrijging van de graad van doctor aan de Technische Universiteit Eindhoven, op gezag van de Rector Magnificus, prof. dr J.H. van Lint,

voor een commissie aangewezen door het College van Dekanen in het openbaar te verdedigen op

vrijdag 28 juni 1991 om 16.00 uur

door

Comelis Wilhelmus Maria Bastlaansen

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Dit proefschrift is goedgekeurd door

de promotoren :

en de copromotor :

prof. dr P.J. Lemstra

prof. dr P. Smith (UCSB)

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Contents

Chapter 1 Introduetion

1.1 Historica} Survey 1.2 Outstanding Problems 1.3 Scope of the Thesis

1.4 References

Part 1 Basic Aspects

Chapter 2 The Influence of Intermolecular Interactions on

the Drawing Behaviour of Polyethylenes.

2.1 Introduetion 2.2 Experimental

2.3

Results 2.4 Discussion 2.5 Conclusions 2.6 References

Chapter 3 On the Nature of Intermolecular Interactions in

Linear Polyethylenes

3.1 Introduetion

3.2

Experimental

3.3

Results

3.4

Discussion

3.5

Condusion

3.6

References 1 1 9 11

14

21 21 22 24 29 30 31 33

33

35

36

44

47

48

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ii

Chapter 4 The Drawing Behaviour of Poly(vinylalcohol)

4.1 Introduetion

4.2

Experimental 4.3 Results 4.4 Discussion 4.5 Conclusions 4.6 Reierences Contents 51 51 52 53 59 61

62 Chapter 5 The Drawing Behaviour of Trans-1,4 Polybutadiene

65

5.1 Introduetion 65

5.2 Experimental 66

5.3 Results 68

5.3.1 Therma1 analysis 68

5.3.2 Tensile drawing behaviour of trans-1,4 polybutadiene 69 5.3.3 Solid state coextrusion of trans-1,4 polybutadiene 70 5.3.4 Post-drawing of coextruded trans-1,4 polybutadiene 73

5.4 Discussion 74

5.5 Conclusions 75

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Contents iii

Part 2 Ultimate Properties

Chapter 6 The Limits of the Tensile Strength of Ultra-Drawn

UHMW-PE Fibres

79

6.1 Introduetion 79

6.2 Literature survey 80

6.2.1 Influence of fibre diameter 80 6.2.2 Influence of chain slippage and chain seission 81 6.2.3 Influence of time-scale and temperature of testing 81

6.3 Experimental 81

6.4 Results 84

6.4.1 Influence of fibre diameter 84 6.4.2 Influence of methyl side groupes 84 6.4.3 Influence of strain rate and temperature 88

6.5 Discussion 91

6.5.1 Influence of fibre diameter 91 6.5.2 Influence of chain slippage and chain seission 93 6.5.3 Influence of time-scale and temperature of testing 94

6.6 Conclusions 95

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iv

Chapter 7 The Theoretical Modulus of Biaxially

Oriented Films

7.1 Introduetion 7.2 Theoretica! aspects

7.2.1 Polyethylene single crystals

7.2.2 Polyethylene single-crystallaminates 7.2.3 Polyethylene single crystal aggregates 7.3 Results

7.4 Discussion and conclusions 7.5 References

Part 3 New Developments

Chapter 8 Orientation Mechanism of Dyes in

Ultra-Drawn Polyethylenes

8.1 Introduetion 8.2 Experimental 8.3 Results 8.4 Discussion 8.5 Conclusions 8.6 References Cootents 97 97

98

99 102 105 107 112 114 119 119 120 121 128 130 131

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Contents

Chapter 9 UHMW-PE Solutions and Gels as a Carrier for

Solid Particles

9.1

Introduetion 9.2 Experimental

9.3 Results and discussion 9.4 Conclusions 9.5 References

Summary

Samenvatting

Dankwoord

133

134

139

140

141

145

V

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Introduetion 1

Chapter 1 Introduetion

1.1 Ristorical Survey

Since the discovery of the first routes for the synthesis of high molecular weight polymers, scientists have been attempting to produce high strength and high modulus structures based on these materials. The prerequisites for actually producing high modulus and high strength materials were already formulated by Carothers and Hi111 in the early 1930's as follows:

"We picture a perfectly oriented fibre as consisring essentially of a single crystal in which the long molecules are in ordered a"ay parallel with the fibre axis".

"Besides being composed of very long molecules, a compound must be capable of crystallizing if it is to farm oriented fibres and orientation is probably necessary for great strength and pliabilily".

A variety of studies was devoted to calculate the theoretica! maximum modulus and strength of synthetic polymers2•13• In genera!, the calculations were performed assuming infinitely long polymerie chains, packed in a perfectly oriented, fully ebain-extended crystal, which is loaded in the chain direction. The theoretica! modulus and strength were calculated based on bending of bond angles and breaking of C-C bands respectively. In the case of polyethylene, the theoretica! modulus2_-S_{and strength}9

"13 were estimated to be 180-400 GPa and 19-35 GPa respectively. These calculations illustrate the potential of synthetic polymers as high modulus and high strength materials and this stimulated research into producing these materials.

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2 Chapter 1

strong and stiff polymerie structures14-30• Experimental routes are nowadays available for producing fibres with Young's moduli and tensile strengtbs which approach their theoretica! values. Two different concepts, based on rigid and flexible macromolecules respectively, have been used for producing these fibres:

Rigid ebains :

lt was discovered by Kwolek et a1.14•15 that, above a critica! concentration in solution, rigid macromolecules, such as PPTA (poly[p-phenylene terephthalamide]), farm liquid crystalline structures which could be spun and oriented in an elongational flow field (figure 1.1 ). After quenching and extraction of the solvent, highly oriented fibres were obtained with excellent mechanica! properties. Nowadays, PPTA fibres are producedon an industrial scale by Dupont (Kevlar) and Akzo (Twaron) possessing Young's moduli and tensile strength's of approximately 90 GPa and 3 GPa respectively.

Air gap

OUjffiCh wat~r - - b a t h

Figure 1.1 : Spinning of PPTA fibres (reproduced with permission from : BunseU A.R.,

Composite Materials Series, part 2, page 271, Elsevier, Amsterdam, 1988)

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Introduetion 3

Flexible ebains :

Intrinsically flexible polymers are also used as a precursor for highly anisotropic, strong and stiff fibres16.:ll>. The development of these fibres took a long and tortuous path. Many different routes were explored and several research groups have contributed to the development. Some results of the more important studies are discussed here. More detailscan be found in several review papers32•34•

The majority of studies on producing strong and stiff structures from flexible polymers were focussed on linear polyethylenes16•30• Extensive studies concerning the deformation mechanism of melt-crystallîzed polyethylenes were performed by PeterJin et al.61ki4. It was shown that, initially, tilting of lamellae in the drawing direction occurs (figure 1.2). Crystalline blocks are formed which reorganize into so-called microfibrils. Further drawing results in shearing of microfibrils and tautening of intra- and inter-microfibrillar tie-molecules.

Figure 1.2 : Drawing mechanism of melt-crystallized polyethylenes (reproduced with permission from : Peter/in A., Coll. Polym. Sci. 364, 265, 1987)

Systematic studies concerning the maximum attainable draw ratio and properties of melt-spun polyethylenes were performed by Ward et al.16-20• It was shown

that the Young's modulus of drawn polyethylene is uniquely related to the draw ratio, independent of, for instance, the molecular weight of the polymer and (below an upper limit) the drawing temperature (figure 1.3).

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4 Chapter 1 60 [J

• .0..

-

el! 50

=...

~ ... til

=

"= _Q ₄₀

e

<1.1 "IX!

=

.;.

30 [J 0 20

r/J'

•

10 0 L---~---~---~----~~ 1 10 20 30 40 draw ratio

H

Figure 1.3 : Young's modulus of melt-crystaUized, drawn polyethylenes

(reproduced with permission from: Capaccio G., Crompton TA., Ward LM., J. Polym. ScL, Phys. Ed., 14, 1641, 1976)

These studies also indicated that the maximum attainable draw ratio of melt-crystallized polyethylenes decreases monotonically with increasing weight average molecular weight (figure 1.4). By optimizing with respect to molecular weight, crystallization history and drawing temperature, fibres with a maximum Young's

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Introduetion

s

modulus and tensile strength of 70 GPa and 1.4 GPa respectively were produced16-20•65•66• The properties of these melt-spun fibres campare favourably with,

for instance, glass fibres. The tensile strength of these melt-spun, ultra-drawn polyethylenes is low compared to, for instance, PPT A fibres and therefore, subsequent studies mainly focussed on increasing the tensile strength.

... ..!... .~

-

=

_"'

~

=

"'

"C 35

25

15

5 1 M,. [kg/mol]

Figure 1.4: Maximum attainable draw ratio of melt-crystallized polyethylenes

(reproduced with permission from: Capaccio G., Crompton TA., Ward LM., J. Polym_ Sci., Phys_ Ed-, 14, 1641, 1976)

An experimental technique to produce polyethylene structures from flowing, supercaoled polyethylene solutions was developed by Penningset al.21_-24_•_{A seed fibre}

was immersed in a Couette flow configuration (figure 1.5) containing a dilute UHMW-PE (Uitra-High-Molecular-Weight Polyethylene,

Mw>HP

kg/mol) solution. The seed fibre was hooked into a polymer layer which formed on the surface of the rotating inner cylinde~1• Subsequently, fibrous structures were grown by withdrawing the seed-fibre.

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6 Chapter l

Figure 1.5: Surface-growth technique (reproduced with permission from: Zwijnenbwg A., Pennings A.J., Coll. Polym. Sci., 254, 868, 1976)

Electron rnicroscopy revealed that the tapes produced consist of agglomerated elernentary fibrils of the "shish-kebab" type (figure 1.6). By optirnizing this technique, the fraction of larnellar overgrowth ("kebab") on the fibres could be reduced. This resulted in UHMW-PE fibrous structures possessing a high Young's modulus (100 GPa) and high tensile strength (3 GPa)21-24• Pennings et al. presented the first

experimental proof that oriented and ebain-extended UHMW-PE structures could be

exceptionally strong and stiff. Unfortunately, the growth rates of the tapes produced were low ( <0.5 m/rnin), and the uniformity along the fibre length, with respect to diameter and strength, was insufficient which are severe limitations for the applicability of this so-called surface growth technique.

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Introduetion 7

Figure 1.6: Shish-kebabs (reproduced with pennission from: Keiler A., Barham P J., Plast.

Rubb. lnst. 6, I, 1979)

lt was shown by Smith and Lemstra27

.30 that, in contrast to melt-crystallized

UHMW-PE, solution-spun/cast UHMW-PE could be drawn in the solid state to high draw ratios, even after complete remaval of solvent from the as-spun fibres. The restrictions encountered by Ward et al.16

·20 in drawing of melt-crystallized UHMW-PE

were circumvented by spinning from semi-dilute solutions. The combination of a high molecular weight (~> 1oJ kg/mol) and a high solid state draw ratio (>30) resulted in

stiff (> 100 GPa) and strong (>3 GPa) polyethylene fibres which could be produced

continuously at high speeds (> 100 m/min.) and with a high uniformity along their length.

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8 Chapter 1

The drawing mechanism of solution-spun UHMW-PE fibres was investigated by Smith, Lemstra and Booij29

• It was found that the maximum attainable draw ratio

of solution-spun UHMW-PE is inversely proportional to the square root of the initial polymer concentration in solution. This experimental observation was interpreted in terrns of an entanglement density dependenee of the maximum attainable draw ratio (figure 1.7). It was assumed that entanglements are "trapped" upon crystallization and limit the draw ratio of UHMW-PE. By crystallization from semi-diJute salution the entanglement density is reduced which enhances the maximum attainable draw ratio (figure 1.7b). Below a critica! concentration (c·) entanglement coupling is absent and consequently the connectivity between individual crystals is lost (figure 1.7c) which results in brittle, non-drawable structures.

a b c

I

'

)

i·

''

::

''

,

)

::

,,

'

J~i'

l •

"

• , , ,

"

(

,,

111

:·

:

:'

I I . ' · I , l t t l t l'

C

l:

:

::::

::

111:11 11• t; ttll I l i l I l i l I l i l I l i l liL I I l i l t i l t l i l t

Figure I. 7 : Entanglement coupling in semi-crystalline polyethylenes

crysta/lized from :a) melt, b) semi-dilute so/ution (c>c*),

I

c) dilure solution (c<c*) (reproduced wirh pennission from: Lemsrra P.J., Kirschbaum R., Polymer, 26, 1372, 1985)

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Introduetion 9

1.2 Outstanding Probieros

In figure 1.8, the present status with respect to the specific modulus and strength of commercial polymerie fibres is depicted. Compared to other synthetic fibres, solution-spun/drawn UHMW-PE fibres (HP-PE) perfarm exceptionallywell. The combination of a high specific modulus and a high specific strength makes these fibres suitable in applications such as helmets, bulletproof vests, ropes, hybrid structural composites etc.32•33•

4,---~---=---,---~

specific stre (N/Tex)

t

JT---+---2 0

•

Nylon

•

Steel 0 50 Boron 100 150 200

Figure 1.8: Speciftc Young's modulus and tensile strengthof commercial

250

ft

bres (reproduced wilh permission from: Lemstra P J., van Aerle NAJ.M., Basliaansen C. WM., Polym. J., 19, 85, 1987)

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10 Chapter 1

Same limitations are encountered, however, in bath the processing and the properties of ultra-drawn UHMW-PE structures. A few of the outstanding probieros were selected and will provide the basis for the thesis.

1. Several attempts have been reported to develop solvent-free routes for the production of high modulus and strength UHMW-PE fibres. Such solvent-free routes have obvious advantages. For instance, it is unnecessary to extract or regenerate large quantities of solvent. Most attempts to develop solvent-free routes were either unsuccessful34_•35 _{or the proposed processes were limited to fabrication of simple}

geometries such as thick fibres, tapes or tube36_·38_•_{The probieros encountered are aften}

related to the drawing behaviour of polyethylenes and more specifically to the influence of melting and recrystallization on the maximum attainable draw ratio of polyethylenes.

2. A rather large research effort was devoted to produce high modulus and strength structures from polar polymers such as poly(vinylalcohol), polyamides etc.3942•

High strength and modulus fibres produced from these materials potentially have some advantages over UHMW-PE fibres such as a higher melting point, less creep under static loading conditions and a higher compressive strength. In most cases, however, the maximum attainable modulus and tensile strength were low in comparison with those of solution-spun/drawn UHMW-PE fibres3942

•

3. Nowadays, drawn UHMW-PE structures are produced with Young's moduli

which closely approach the theoretica! modulus43

•44 (Eexp=[0.6-0.9]E_1h).In contast, the

tensile strengthof UHMW-PE fibres is still rather low45•46 compared to the theoretica!

values (a₁,exp=[O.l-0.3]a_1,1h).A variety of studies has been performed concerning the fracture mechanism of UHMW-PE fibres47"51• The results of these studies are aften

contradictory because different parameters are used in the rnadelling of the tensile strength. Consequently, the origin of the limited tensile strength of solution-spun, drawn UHMW-PE fibres is still a matter of debate.

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Introduetion 11

4. Equi-biaxial drawing is aften used to enhance the mechanica] properties of polymerie films. At first sight, it seems obvious to use dissalution routes to produce biaxially drawn, high modulus UHMW-PE films. However, compared with uniaxial drawing, the maximum attainable Young's modulus (

<

10 GPa) and tensile strength

( <

1 GPa) are lo~2•53• Moreover, reliable estimates with respect to the theoretica! modulus and strength of equi-biaxially drawn polymers are not available until now, despite their obvious practical relevance.

5. Ultra-drawing of semi-crystalline polymers has been used mainly to produce high modulus and strength structures. Both high draw ratio's and high orientation levels can also be used to enhance other properties of polymerie systems57-59• However, little attention has been devoted to exploring this potentially interesting research area.

1.3 Scope of the Thesis

A variety of subjects concerning the processing, drawing behaviour and properties of semi-crystalline polymers is discussed. Three main topics are to be distinguished : Basic Aspects (part 1 ); Ultimate Properties (part 2); New Developments (part 3).

Part 1 : Basic aspects of the drawing behaviour and properties of semi-crystalline polymers

In chapters 2-5, the drawing behaviour of some semi-crystal1ine polymers is discussed. In chapter 2, a sealing law for the maximum attainable draw ratio of linear polyethylenes is derived from experimental data. Subsequently, it is attempted to further identify the physical nature of the constraints which limit the maximum attainable draw ratio of polyethylenes ( chapter 3). The results are discussed in relation to solvent-free routes for the production of high modulus and strength structures.

In chapter 4, the drawing behaviour and properties of poly(vinylalcohol) is discussed. Compared to linear polyethylenes, additional intermolecular interactions, in the form of hydragen honds, are present in poly(vinylalcohol), which reduce their

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12

Chapter 1

drawability.

In chapter 5, the drawing behaviour of trans-1,4 polybutadiene is investigated. Trans-1,4 polybutadiene is a highly unsaturated polymer and can potentially be used to produce crosslinked high modulus and high strength structures with, for instance, a Iow level of creep under static loading condîtions.

Part 2 : Ultimate properties

In the second part of the thesis ( chapters 6, 7), the limits of the properties (tensile strength, theoretical equi-biaxial modulus) of drawn flexible polymers are explored.

In chapter 7, the fracture behaviour of UHMW-PE fibres is investigated. Based on a series of experiments, it is attempted to distinguish between several roodels which are currently used to describe the fracture behaviour of UHMW-PE fibres.

In chapter 8, a model is presented to calculate the theoretica! modulus of equi-biaxially oriented structures. Based on these calculations, the prerequisites for producing high modulus films are derived.

Part 3: New Developments

Same new applications of ultra-drawn semi-crystalline polymers are briefly discussed in an attempt to generate other favourable properties than a high Young's modulus and tensile strength via solid state drawing.

In chapter 8, the orientation mechanism of dichroic dyes in drawn polyethylenes is presented. It is shown that high quality polarizers can be produced from dye-containing, uniaxially drawn, polyethylene films.

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Introduetion 13

In chapter 9, it is shown that salution-cast UHMW-PE films can be used as a carrier for large quantities of solid particles. Moreover, biaxially drawn UHMW-PE films, containing ceramic particles, appear to be excellent precursors for (thin) ceramic films.

This thesis is based on a collection of papers which have been publisbed in, or have been submitted to, various joumals67

-78• Furthermore, the author bas contributed to

some papers on related subjects which are nat presented in this thesis 79 -80.

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14 Chapter 1

1.4 References

1) Carothers W.H., Hill J.W., J. Am. Chem. Soc., 54, 1579, (1932) 2) Treloar L.R.G., Polymer, 1, 95, (1960)

3) Odajima A., Maeda T., J. Polym. Sci., Part C, 15, 55, (1966) 4) Wobser G., Blasaenberg S., Coll. Polym., Sci., 241, 985, (1970) 5) PerepeHein K.E., Angew. Makrom. Chem., 22, 181, (1971) 6) Boudraux D.S., J. Polym. Sci., Phys. Ed.,ll, 1285, (1973)

7) Tashiro

K.,

Kobayashi M., Tadokoro H., Macromolecules, 11, 914, (1978) 8) Crist B., Ratner M.A., Brower AL., Sabin J.R., J. Appl. Phys., 50, 6047, (1979) 9) Zhurkov S.N., Intern. J. Fract. Mech., 1, 311, (1965)

10) Zhimov N.l., Koryak-Doronenko Y.G., Hartelev G.M., Polym. Sci. USSR, 11, 1396, (1969)

11) Mark H.F., "Polymer Science and Materials", Ed. Tobolski A.V. and Mark H.F., Wiley Interscience, New York, 231, (1971)

12) He T., Polymer, 27, 253, (1986)

13) Kelly A., Macmillan N.H., "Strong Solids", 3rd ed., Claredon Press, Oxford, (1986) 14) Kwolek S.L., USP 3, 671, 542 (1972)

15) Kwolek S.L., Morgan P.W., Schaefgen P.W., Gulrich L.W., Macromolecules, 10, 1390, (1977)

16) Capaccio G., Ward I.M., Nature (Phys. Sci), 243, 143, (1973) 17) Capaccio G., Ward I.M., Polymer, 15, 233, (1974)

18) Capaccio G., Ward I.M., Polymer, 16, 239, (1975)

19) Capaccio G., Ward I.M., Polym. Eng. Sci., 15, 219, (1975)

20) Capaccio G., Crompton T.A., Ward I.M., J. Polym. Sci., Phys. Ed., 14, 1641, (1976) 21) Pennings AJ., Van der Mark J.M.A.A., Booij H.C., Coll. Polym. Sci. 99, 236,

(1970)

22) Zwijnenburg A., Pennings AJ., Coll. Polym. Sci., 452, 253, (1975) 23) Pennings A.J., J. Polym. Sci., Polym. Symp., 55, 59, (1977)

24) Zwijnenburg A., PhD Thesis, University of Groningen, The Netherlands, (1978) 25) Gibson AG., Ward I.M., Cole B.N., Parsons B., J. Mater. Sci., 1193, 9, (1974) 26) Perkins W.G., Capiaty N.G., Porter R.S., Polym. Eng. Sci., 3, 16, (1976) ·

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Introduetion 15

27) Smith P., Lemstra P.J., Polymer, 21, 1341, (1980)

28) Smith P., Lemstra P.J., Kalb B., Pennings AJ., Polym. Bull., 1, 733, (1979) 29) Smith P., Lemstra P.J., Booij H.C., J. Polym. Sci., Phys. Ed., 19, 877, (1981) 30) Smith P., Lemstra P.J., Coll. Polym. Sci. 258, 891, (1980)

31) Barham P.J., Keiler A, 20, 2281, (1985)

32) Lemstra P.J., Kirschbaum R., Ohta T., Yasuda H., "Oevelopments in Oriented Polymers-2", Ed. Ward I.M., Elsevier Applied Science Publ., New York, chapter 2, (1987)

33) Kirschbaum R., Yasuda H., Van Gorp E.H.M., Chemiefasern!fextielind. T134-T139, (1986)

34) Lemstra P.J., Kirschbaum R., Polymer, 26, 1372, (1985)

35) Bastiaansen C.W.M., Froehling P., Pijpers AJ., Lemstra P.J., "Integration of Polymer Science and Technology", Ed. Kieintjens L., Lemstra P.J., Elsevier Applied Science Publ., 508, (1985)

36) Kanamoto T., Tsurata A, Tanaka M., Porter R.S., Polym. J., 15, 327, (1983) 37) Kanamoto T., Tsurata A, Tanaka M., Takeda M., Polym. J., 16, 75, (1984) 38) Smith P., Chanzy H.O., Rotzinger B.P., Polym. Comm., 257, 26, (1985) 39) Prevorsek O.C., Kwon Y.O., Sharma R.K., J. Mater. Sci., 12, 2310, (1970) 40) Acirno A, Mantia F.P., Polzzotti G., Cifferi A, J. Polym. Sci., Phys. Ed., 17, 1903,

(1970)

41) Gogolewski S., Pennings AJ., Polymer, 26, 1394, (1985) 42) Perepelkin K.E., Ang. Macrom. Chem., 22, 181, (1972)

43) Matsuo M., Inoue K., Abumiya N., Sen-i Gakkaishi, 40, T275, (1984) 44) Kunugi T., Oomori S., Mikami S., Polym., 29, 814, (1988)

45) Hoogsteen W., PhO Thesis, University of Groningen, The Netherlands, (1989) 46) Hoogsteen W., ten Brinke G., Pennings AJ., Coll. Polym. Sci., 266, 1003, (1988) 47) Smith P., Lemstra P.J., J. Polym. Sci., Phys. Ed., 19, 1007, (1980)

48) Smith P., Lemstra P.J., J. Polym. Sci., Phys. Ed., 20, 2229, (1982) 49) Smook J., Hamersma W., Pennings AJ., J. Mater. Sci., 19, 1359, (1984) 50) Wagner H.O., Steenbakker L.W., Phil. Mag. Lett., 59, 77, (1989) 51) Termonia Y., Meakin P., Smith P., Macromolecules, 18, 2246, (1985)

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16

Chapter 1

52) Gerrits N., PhD Thesis, Eindhoven University of Technology, the Netherlands, (1990)

53) Minami S., Itiyama K., Polym. Prepr. Am. Chem. Soc., Div. Polym. Chem., 26, 2, 245, (1985)

54) Clements J., Ward I.M., Polymer, 23, 935, (1982)

55) Smook J., PhD Thesis, University of Groningen, The Netherlands, (1984) 56) Lemstra P., Kirschbaum R., Polymer (conf. issue), 26, 1372, (1985)

57) Greig D., "Developments in Oriented Polymers-1", Ed. Ward I.M., Appl. Sci. Publ. London, (1982)

58) Chiang J.C., Smith P., Heeger A.J., Wudl F., Polym. Comm., 29, 161, (1988) 59) Poulart B., Chielens J.C., Vandenbende C., Issi J.P., Legras R., Polym. Comm., 31,

148, 1990

60) PeterJin A., J. Polym. Sci., C 9, 61, (1%5) 61) PeterJin A., J. Mater. Sci., 6, 490, (1971) 62) PeterJin A., Coll. Polym. Sci., 253, 809, (1975) 63) PeterJin A., J. Appl. Phys., 48, 4099, (1977) 64) PeterJin A., Coll. Polym. Sci., 265, 357, (1987) 65) Wu T., Black H., USP 42228118, (1980)

66) Wu T., Black H., Polym. Eng. Sci., 19, 163, (1979)

67) Bastiaansen C.W.M., J. Polym. Sci., Phys. Ed., 23, 2242, (1990)

68) Schellekeos R., Bastiaansen C.W.M., J. Appl. Polym. Sci., accepted for publication 69) Bastiaansen C.W.M., Meijer H.E.H., Lemstra P.J., Polymer, 31, 1435, (1990) 70) Bastiaansen C.W.M., Froehling P., Pijpers A.J., Lemstra P.J., "lntegration of

Polymer Science and Technology". Ed. Kieintjens L, Lemstra P.J., Elsevier Applied Science, London, p. 508, (1985)

71) Lemstra P.J., Bastiaansen C.W.M., Meijer H.E.H., Ang. Makrom. Chem, 145/146, 343, (1986)

72) Lemstra P.J., Van Aerle N.A.J.M., Bastiaansen C.W.M., Polym. J., 19, 85, (1985) 73) Leblans P.J.R., Bastiaansen C.W.M., Macromolecules, 22, 3312, (1989)

74) Van Aerle N.A.J.M., Lemstra P.J., Kanamoto T., Bastiaansen C.W.M., Polymer, 32, 34, (1990)

(27)

Introduetion

17

76) Bastiaansen C.W.M., Grooters G.P., accepted for pubHeation by Polymer 77) Bastiaansen C.W.M., Leblans P.J.R., Smith P., Macromolecules, 23, 236, (1990) 78) Bastiaansen C.W.M., Leblans P.J.R., Smith P., accepted for pubHeation by

Macromolecules

79) Bastiaansen C.W.M., Lemstra P.J., Makrom. Chem., Makrom. Symp., 23, 73, (1989) 80) Motamedi F., Bastiaansen C.W.M., Smith P., submitted for pubHeation to

(28)

(29)

Part 1 Basic Aspects of the Drawing Behaviour of

Semi-Crystalline Polymers in the Solid State

(30)

(31)

Intennolecular interactions

Chapter 2 The lnfluence of Intermolecular Interactions on the

Drawing Behaviour of Polyethylenes

2.1 Introduetion

21

Extensive studies have been performed concerning the drawing behaviom of linear po1yethylenes1•10• The prime objective of these studies was togeneratea high

degree of chain orientation and extension via drawing at elevated temperatures in a temperature range close to but below the melting temperature.

Systematic studies conceming the drawing behaviour of melt-crystallized polyethylenes were performed by Capaccio and Ward1·5• It was shown that, at optimum

conditions with respect to crystallization history and drawing temperature, high maximum attainable draw ratios could he obtained using commercial Iinear polyethylenes (Mw-102 _{kg/mol). However, the maximum attainable draw ratio of}

melt-crystallized polyethylenes decreases monotonically with increasing molecular weight1·5 (see figure 1.4, chapter 1).

The drawing behaviour and properties of salution-cast polyethylenes were extensively studied by Smith, Lemstra and Booit"10• The maximum attainable draw

ratio of High-Molecular-Weight Polyethylenes (HMW-PE, Mw>5x10Z kg/mol) and Ultra-High-Molecular-Weight Polyethylenes (UHMW-PE, Mw>l@ kg/mol) could be enhanced by spinning and casting from semi-dilute solutions. Moreover, these studies also indicated that the excellent drawing characteristics of salution-cast UHMW-PE films were preserved even after complete remaval of solvent from the films.

This chapter is reproduced, in part, from :

1. Basliaansen C. W.M, J. Polym. Sci., Phys. Ed., 23, 2242,, (1990)

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22 Chapter 2

The maximum attainable solid state draw ratio (lmax) of polyethylenes is influenced by bath the molecular weight (Mw) of the polymer1•5 and the polymer concentration ( c) in solution<HO. In this chapter, additional experimental results are presented conceming the influence of both these parameters, Mw and c, on the draw ratio of polyethylenes and a universa! sealing law is derived.

2.2 Experimental

The molecular weight (Mw) and polydispersity (Q= Mw!Mn) of the polyethylene grades used in this study are listed in table 2.1. These polyethylene grades were kindly supplied by DSM (The Netherlands, grade a), Hoechst/Ruhrchemie (Germany, grade b,d), Mitsui (Japan, grade c) and Himont (USA, grade e)

Table 2.1: Weight average molecular weight and polydispersity of

the

polyethylene

grades used in this study

code a b c d e

Mw

[kg/mol] 1.2xlo2 4.5x102 8.0x102 l.Sxlo-' 4.5xlo-'

Mw/Mn

[-] 8 8 11 9 5

Prior to the dissalution procedure, 1 % w/w di-butyl-p-cresol was added to the polymers to prevent degradative oxidation. Xylene was added to the polymers and the polymer-solvent mixtures were degassed at room temperature. Subsequently, the polymers were dissolved in xylene at 130 °C. During the dissalution procedure, the solutions were stirred until the so-called Weissenberg-effect appeared. After dissalution

(33)

Intermolecular interactions 23

occurred, which took approximately 2 hours, the solutions were quenched to room temperature. After complete evaporation of solvent at ambient conditions, solvent-free films were obtained (table 2.2).

Table 2.2: Weight average molecular weight

(Mw)

and inititial polymer

concentradon in salution ( c) of the polyethylene films

code 1 2 3 4

5

6 7 8 9 10

Mw

[kglmol] 1.2x1Q2 4.5x1Q2 S.Oxl(f &.Oxl(f l.Sx103 1.5xld3 1.5x103 4.8xld3 4.8xld3 4.8xld3 c [% w/w] 40*

1o·

10

25.

10 2.5

o.s·

10

25

0.3·

• Polyethylene films cast from solutions with a polymer concentradon just above

the critica/ concentradon to obtain coherent films

The critical concentratien to obtain coherent, ultra-drawable films (c·) was determined by testing whether the films were drawable and coherent at a temperature of 100 °C7•11• Brittle, incoherent films were excluded from this study. Polyethylene films

cast from solutions with a polymer concentration just above c • are marked with an asterisk in table 2.2.

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24 Chapter 2

Melt-crystallized polyethylene films were prepared by compression-mouldingof pellets (grade a) or reactor powders (grade b-d) at 180 °C for 30 minutes and subsequently quenching to room temperature.

Tapes were cut from the salution-cast and melt-crystallized films. Solid state drawing of these polyethylene tapes was performed on thermostatically controlled hot shoes at a constant (Hencky) strain rate of approximately 0.05 s·1_•_{The draw ratio was}

determined by measuring the displacement of ink-marks.

2.3 Results

In the semi-diJute regime, the maximum attainable draw ratio of polyethylenes increases with decreasing initia} polymer concentratien in solution6-10• At low initia!

polymer concentrations, in the dilute regime, brittie and non-drawable films are obtained6_-10_{_}_{In other words, an optimum with respect to maximum attainable draw}

ratio exists at an initia! polymer concentration in salution which slightly exceeds the critica! concentration to obtain coherent, ultra-drawable salution-cast films (c·). In accordance with previous studies14, it is found that this critica! concentration decreases

with increasing weight average molecular weight (see experimental section and table 2-2). Moreover, the experimental results in figure 2.1 show that, at optimized conditions with respect to polymer concentration in salution (c just above c·), high maximum attainable draw ratios can be obtained with high molecular weight polyethylenes.

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Intermolecular interactions

25

The maximum attainable draw ratio of films, at a fixed polymer concentration in solution, is plotted in figure

2.1

as a function of weight average molecular weight.

lt is shown that the attainable draw ratio, at a ftxed polymer concentration, decreases monotonically with increasing molecular weight.

,....,

• ...,

.$!

-

f

ll';

f

"= 70 60 50 40 30 20 10 1

~~.~--~2---+4-+6-8~~~----2---4--~8-.+

M,. [kg/mol]

Figure 2.1: Maximum attainable draw ratio at 100 °C of polyethylene films. (o) melt-crystallized PE, (x) c=10% w/v, (•) c=2.5% w/ll, ( •) data with c just above c •

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26 Chapter 2

The maximum attainable draw ratio of salution-cast films, with a polymer concentration of 10 % w/v, is plotted in figure 2.2. In these experiments, UHMW-PE tapes we re either drawn isothermally at 100

°

e

or subjected to multi-stage drawing in an increasing temperature gradient from 100

oe

to 130

oe

and 100

oe

to 150

oe.

,...., I ... 40 20 10

~~~~~~~~~

}()l 2 4 6 8

lo:'

2 4 6 8

loot

M..

[qimol]

Figure 2.2 : Maximum attainable draw ratio of solution-cast UHMW-PE films (c=10% w/v). (o) isothennal drawingat 100 °C, (+) multistage drawing 100-130 °C, (x) multistage drawing 100-150

oe

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Intennolecular interactions

27

In figure 2.3, the maximum draw ratio of melt-crystallized polyethylenes is plotted (on a double logarithmic scale) as a function of the weight average molecular weight.Experimental data from previous studies by Ward et al.3 _{and data obtained in}

the present study are shown. Complete sets of experimental data were shifted vertically, over an arbitrary distance, to correct for differences in drawing temperature and strain rate during drawing. The vertical shift of the experimental data does not influence the power-law exponent in the relation between the maximum draw ratio and the molecular weight. A linear re]ationship is observed with a slope of -0.5 +0.05 which indicates that the draw ratio scales with the molecular weight as follows:

-

I ...

.s

-

f!

!1:

f!

"C:: -0.5±0.05 Àmax-

Mw

2

)<,

'x

101 8 6 4 2

1 ~4-~--~

1()2 2 4 ₆

₈₁₀₃

₂ M,. [kg/mol] 4 ₆

_810'

(1)

Figure 2.3 : Maximum attainable draw ratio of melt-crystallized polyethylenes.

(o) data from figure 2.1., (x) data from reference 3 (experimental data were shifted vertically to correct for differences in drawing temperature)

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28 Chapter 2

In figure 2.4, the maximum attainable draw ratio of melt·crystallized and salution-cast polyethylenes is plotted as a function of the Bueche parameter ( c.Mw).

8 6

~to

4

~

,..., 0 \ 0

• ....

. s

-

₂

~x

=

...

~ ~+

f

'+-,~ "C

_'\..o+

10 _~

8

x

6

+

4 0 2 1

~~+-~----~~

₂ ₄ _{6 8}

₁₀₁

₂ ₄ _{6 8}

₁₍₎₃

₂ ₄ _{6 8}

_10'

c.M.,

Figure 2.4 : Maximum attainable draw ratio of melt-crystallized and salution-cast polyethylenes. (o) data from figures 2.1, 2.2 and 2.3, (x) data from re[ erenee 3, ( +) data from reference 7. ( experimental data were shited vertically to co"ect for differences in drawing temperature)

(39)

lntermolecular interactions

29

Experimental data from previous studies3•7 are also shown in figure 2.4. Again

complete sets of data were shifted vertically to correct for the influence of drawing temperature on the maximum attainable draw ratio. A linear relationship with slope of -0.5±0.05 is observed, which shows that the maximum attainable draw ratio scales with the Bueche parameter as:

-0.5±0.05

Àmax - (

c.M_")

2.4 Discussion

(2)

The influence of intermolecular interactions (entanglements) on the rheological behaviour of polymer melts and solutions has been investigated previously1_z..17_•_For

polymer solutions, in genera!, a distinction is made between the rheological behaviour of solutions below and above a critica! concentradon ( c

*)

which corresponds to the onset of entanglement coupling in solution. The rheological properties of concentrated solutions (c>c') are aften correlated with the so-called Bueche parameter (c.Mw). This parameter is proportional to the number of intermolecular cantacts per molecule and dominates the rheological properties of concentrated solutions. For instance, the zero-shear viscosity of concentrated solutions is proportional to c.Mw to the power 3.4, if corrections are made for the concentration dependenee of the friction factor15

•

In figure 2.4, it was show that the maximum attainable draw ratio of

polyethylenes is inversely proportional to the square root of c.Mw, over the entire range of experimental conditions investigated. Apparently, this Bueche parameter not only influences the rheological properties of melts and solutions12

"17, but also dominates the

solid state drawing behaviour of polyethylenes. Therefore, it is concluded that intermolecular interactlans in polymer melts and solutions, determine the maximum attainable draw ratio of polyethylenes.

In semi-crystalline polar polymers, such as poly(vinylalcohol), additional intermolecular interaction are present in the crystal Jattice in the form of rather strong hydragen bands. The influence of these hydragen bands on the drawing behaviour of poly(vinylaJcohol) will be dicussed in chapter 4.

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30 Chapter 2

2.5 Conclusions

It was shown that the maximum attainable draw ratio of melt-crystallized and salution-cast polyethylenes is inversely proportional to square root of the Bueche parameter ( c.M.v). This experimental observation indicates that the maximum draw ratio of polyethylenes is determined by the number of intermolecular interactions per molecule. The physical nature of these intermolecular interactions is discussed in the next chapter.

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Intermolecular interactions 31

2.6 References

1) Capaccio G., Ward I.M., Nature (Phys. Sci.), 243, 143, (1973) 2) Capaccio G., Ward I.M., Polymer, 15, 233, (1974)

3) Capaccio G., Ward I.M., Polymer, 16, 239, (1975) 4) Capaccio G., Ward I.M., Polym. Eng. Sci., 15, 219, (1975)

5) Capaccio G., Crompton T.A., Ward I.M., J. Poly. Sci., Phys. Ed., 14, 1641, (1976) 6) Smith P., Lemstra P.J., DSM/Stamicarbon, U.S. Pat. 4, 344, 908; 4, 422, 933;

4, 430, 383; 4, 436, 689, (1978)

7) Smith P., Lemstra P.J., Booij H.C., J. Polym. Sci., Phys. Ed., 19, 877, (1981) 8) Smith P., Lemstra P.J., Polymer, 21, 1341, (1980)

9) Smith P., Lemstra P.J., Kalb B., Pennings A.J., Polym. Bull., 1, 733, (1979) 10) Smith P., Lemstra P.J., Coll. Polym. Sci., 258, 891, (1980)

11) Sawatari C., Tomoko 0., Matsuo M., Polym. J., 18, 741, (1986) 12) Gupta D., Forsman W.C., Macromolecules, 3, 304, (1969) 13) Allen V. R., Fox T.G., J. Chem. Phys., 41, 337, (1964) 14) Berry G. C., J. Phys. Chem., 70, 1194, (1966)

15) Berry G. C., Fox T.G., Adv. Polym. Sci., 5, 261, (1968) 16) Graessley W.W., Adv. Polym. Sci., 16, 4, (1974)

17) Ferry J. D., "Viscoelastic Properties of Polymers", 2nd ed., Wiley, New York, (1970)

18) Smith P., Lemstra P.J., J. Polym. Sci., Polym. Phys. Ed., 19, 1007, (1981)

19) Smith P., Lemstra P.J., Pijpers J. P. L., J. Polym. Sci., Phys. Ed.,

20,

2229, (1982) 20) Bastiaansen C.W.M., Meijer H.E.H., Lemstra P.J., Polymer, 31, 1435, (1990) 21) Schellekens R., Bastiaansen C.W.M., accepted for pubHeation by J. Appl. Polym.

Sci.,

22) Schellekens R., Bastiaansen C.W.M., submitted for pubHeation to J. Appl. Polym. Sc i.

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(43)

Intermolecular lnteractions 33

Chapter 3 On the Nature of Intermolecular Interactions in Linear

Polyethylenes

3.1 Introduetion

In chapter 2, it was shown that the maximum attainable draw ratio of polyethylenes in the solid state is inversely propotional to the square root of the Bueche parameter ( c.Mw, where c is the initial polymer concentration in salution and Mw is the molecularweight). This experimental observation indicates that the maximum draw ratio depends on the number of intermolecular interactions per molecule.

In the past, several attempts have been made to identify the nature of intermolecular interactions in polymer melts and solids1•2• In figure 3.1, a schematic

drawing of an intermolecular interaction in polymer melts is shown. Intermolecular interactions are represented by topological constraints ( entanglements) which are visualized as four strands leading away from a contact1•

This chapter is reproduced, in

part,

from :

1. Bastlaansen C. W.M., Meijer H.E.H., Lemstra P.J., Polymer, 31, 1435, (1990) 2. Bastlaansen C. W.M., Froehling P., Pijpers A.J., Lemstra P.J., "Integratlon of Polymer

Science and Technology'" Elsevier Appl. Sci., London, p. 508, 1985

3. Lemstra P.J., Bastlaansen C. W.M., Meijer H.E.H., Ang. Makrom. Chem., 145!146, 343, 1986

4. Lemstra P.J., Van Aerle NA.J.M., Basliaansen C. W.M., Polym. J., 19, 85, (1987) 5. Leblans P.J.R., Bastiaansen C.W.M., Macrom., 22, 3312, (1989)

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34 Chapter 3

Figure 3.1: Schematic drawing of an entanglement (according to Ferry in reference 1)

Further information with respect to the physical nature of entanglement coupling can potentially be obtained from experimental studies on the timescale of memory effects, upon heating in the melt, of polymers with a reduced entanglement density. If the assumed physical nature of entanglements in figure 3.1 is correct than diffusion of macromolecules along their own contour length (i.e. reptation) is a prerequisite for "re-entangling" of disentangled polymers.

A model concerning the molecular mobility of polymer ebains in the melt was actvaneed by De Gennes3

•4• A constitutive equation, basedon this model, was derived

by Doi and Edwards35• In these models, it is assumed that the longest relaxation times

in polymer melts correspond to reptative motion of complete chains. The relaxation times conesponding to reptation we re determined experimentally by Klein and Bricoe5•

These relaxation times are extremely long in high molecular weight or branched polymers5•

In previous studies long time memory-effects, related to the entanglement density of polymers, were indeed observed. Por instance, repeatedly extruded, "shear-refined" Low Density Polyethylene (LDPE) exhibits a drop in melt-viscosity which was attributed to "disentangling" during repeated processing. The viscosity drop could be

reversed by heating in the melt for a prolonged period of time (t> > HY seconds at T>160 °C).

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Intermolecular Interactions 35

In this study experimental results on memory effects in solution-cast LDPE and UHMW -PE are presented. The life time of memory effects in the melt of salution-cast polyethylenes is investigated and related to macromolecular mobility and the physical origin of entanglement coupling.

Low

Density Polyethylene (LDPE) and UHMW-PE are chosen as model matenals because of their long relaxation times in the melt. eonsequently, long term memory effectscan potentially be observed in these materials. Based on the experimental data, it is attempted to gain a better understanding of the role of intermolecular interactions on the drawing behaviour of linear polyethylenes.

3.2 Experimental

The polyethylene grades used are Stamylan 2800 (LDPE,

Mw=1.5xlo2

kg/mol) and Hifax 1900 (UHMW-PE,

Mw=5xHf

kg/mol).

To the polymers, 1 % w/w di-butyl-p-cresol was added to prevent degradative oxidation. Subsequently, polyethylene solutions in xylene with a nomina! polymer concentration of 1 % w/v were prepared at 130 oe. After complete dissalution occurred, which took approximately 2 hours, the solutions were cast and quenched to room temperature. Subsequently the solvent was evaporated at ambient conditions. In

the case of LDPE, brittle, incoherent and porous films were obtained. These films were compression moulded in the solid state at 75 oe in order to obtain coherent structures. Salution-cast films, wrapped in aluminium foil, were melted and recrystallized by immersing in a preheated silicone oil bath. After a preset time the films were quenched to room temperature.

Melt-crystallized LDPE and UHMW-PE films were prepared by compression-moulding of respectively pellets and reactor powders at 180 oe for 30 minutes. It is assumed that these melt-crystallized LDPE and UHMW-PE samples possess equilibrium viscoelastic properties in the solid state and melt.

Stress-straio measurements in the solid state were performed on a Zwick Tensile Tester equipped with a thermostatically controlled oven. Dumbbell-shaped samples with an original length of 10 mm were drawn at a constant crosshead speed of 100 mm.min·1•

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36

Chapter 3

performed on a Rheometrics Mechanica! Spectrometer, type RMS 7200, equipped with a paraHel-plate system. Before the measurements were started, the samples were heated for approximately 3 minutes to obtain thermal equilibrium.

Uniaxial drawing experiments in the melt were performed on a Göttfert Rheostrain at a constant strain rate of 0.1 s·1

• A detailed description of the apparatus is given in references 33 and 34. Dumbbell-shaped samples with an originallength of

20 mm were used. Prior to the measurements, the samples were heated for approximately 1 minute.

3.3 Results

Solution-cast LDPE

In figure 3.2 nomina! stress-strain curves, recorded in the solid state, of salution-cast (S) and melt-crystallized (M) LDPE are shown. The salution-cast sample was crystallized from a salution with a polymer concentration below the critical concentradon for coil overlap. Consequently, the connectivity between individual crystals after quenching to room temperature is lost, which results in a low maximum attainable draw ratio.

An

identical situation is obtained in the case of UHMW-PE, if films are cast at a much lower concentradon ( chapters 1,2)6•

The solid state drawing behaviour of recrystallized, salution-cast LDPE films (S*) is similar to the melt-crystallized reference sample (M) (figure 3.2). Apparently, the dissalution history of salution-cast LDPE films, with respect to solid state drawability, is lost after heating in the melt. The surprising observation in these experiments is the short time scale ( < 60 seconds) needed to restare equilibrium properties in the solid state ( campare with shear-refined LDPE).

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Intermolecular Interactions

37

10 M S* ,..., al ~ ~ ... (I) 5 (I)

e

....

_(I) draw ratio [ ·]

Figure 3.2: Nomina/ stress-strain curves of LDPE at 75 °C.

(M) melt-crystallized WPE,

(S) solution-cast LDPE, (S•) solution-cast, recrystaUized WPE

(1

min at 140 °C)

In figures 3.3 and 3.4, dynamic measurements and elongational flow measurements of salution-cast (S) and melt-crystallized (M) LDPE samples are shown respectively .. The dynamic mechanical and elongational flow measurements were started after heating for respectively 3 minutes and 1 minute to 120 °C to allow for thermal equilibrium. In bath cases no differences are observed in rheological behaviour between the salution-cast and melt-crystallized samples.

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38 Chapter 3 60 6 ...., 41

:

-

_OAI 50

...

_...

=

t-'

=

u <IJ

~x

x~

.i

<IJ

=

-

40

x /

5

~x

30

/

~x

x

'o.

/

20 4 10 ·1 0 1 2 ·1 0 1 2 loge.> [rad/s]

Figure 3.3: Loss angle and dynamic modulus

Gd

at 120 °C of melt-crystaUized

(x)

and solution-cast (o) LDPE.The measurements were started aftera hearing

time of 3 minutes at 120 °C.

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39

time [s)

Figure 3.4 : Stress build-up, at a strain rate of 0.1 s-1 _{and a temperature of 120 °~}_of

melt-crystallized ( +) and salution-cast (o) WPE. Ihe measurements were started aftera healing time of 1 minute at 120 °C.

In tigure 3.5, the results of dynamic mechanica} measurements of LDPE are shown. The phase angle and dynamic modulus are measured in a temperature range from 120 to 210

oe

for angular frequencies ranging from

w-

1 _{to 10Z rad/s. A master}

curve was constructed by first shifting the phase angle along the horizontal axis, then shifting the modulus curves over the same distance along the frequency axis and finally effecting superposition of the modulus curves via a small shift along the modulus axis7•

From this master curve the relaxation-time spectrum can be derived. The frequency at which the phase angle starts to deviate from 9(f is of the order of 104 rad/s at

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40 Cbapter 3

190 °C. This deviation occurs when the product of angular frequency. and longest relaxation times in the material is approximately unity. In other words, the experimental results in figure 3.7 indicate that relaxation times over 104 seconds are present in the LDPE melts at 190 °C.

'

80

'

_' 5

,....,

_t

41

x\

-

_~ ...

=

..

=

c

~ ~+ ₄ I:JG ~ _e

.s

60

\

I

-x

I

3 x 40 _I I I ₂ I I I 20 I 1 0 0 -4 -2 0 2 4-4 ·2 0 2 4 log <..> [rad/s]

Figure 3.5: Master curve at 190

oe

of the phase angle and dynamic modulus of WPE. Measuring temperatures: (•) T=l20

°C: (•)

T=150

°C: (")

T=l70 °Ç,

(51)

lntennolecular Intemctions

41

Solution-cast Ultm-High-Molecular-Weight Polyethylene

Experimental results on memory effects in salution-cast UHMW-PE have already been reported previouslyll-10• However, in these studies a relatively low molecular weight sample (Mw= 1.5xlo3 kg/mol) with a high polydispersity (Q=

M.JMn

=

10-15) was used, which causes problems in interpreting the experimental data. Therefore these experiments were repeated using a UHMW-PE grade with a higher weight average molecular weight (Mw=4.8x1<P kg/mol) and a lower polydispersity (0=3-5).

Stress-strain measurements in the solid state and dynamic mechanical measurements in the melt of this particular UHMW -PB grade are shown in figures 3.6 and 3.7 respectively. 40 ,...,

=

1:1.. ~ _M ... til ₃₀

s·

til

e

..

(I} 20 10

s

0~---.---r---~---.----1 10 20 30 40 draw ratio [-]

Figure 3.6 : Nomina/ stress-strain curves of UHMW-PE at 90 °C.

(M) melt-crystallized UHMW-PE, (S) salution-cast UHMW-PE, (S•) solution-cast, recrystallized UHMW-PE (1 min at 150 °C}

(52)

42 Cbapter 3

In the dynamic mechanica] measurements 3 minutes were allowed in order to reach thermal equilibrium. Once more it is observed that the dissalution history of UHMW-PE with respect to solid state and melt properties is lost almast instantaneously upon heating into the melt. Unfortunately, it is impossible to perfarm uniaxial drawing experiments in the true melt (T> 160 °C) because the UHMW-PE samples fail at small uniaxial deformations10•

60 107 ~ ';i,

=

50

=

,...

<1.1

=

<1.1 ~

.s

...

40

~

106 ~ CIQ e

-30 20 105 10

~

0 104 -1 0 1 2 -1 0 1 2 loge.>

[radls]

Figure 3.7: Loss angle and dynamic modulus Gd at 160

oe

ofmelt-crystallized (o) and salution-cast ( •) UHMW-PE. The measurements were started after healing to 160

oe

for 3 minutes.

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In figure 3.8, the stress relaxation modulus of UHMW-PE is shown. The measurements were stopped after approximately 3 hours. This figure shows that relaxation times over 104 _{seconds are present, even at 180 °C, in this particular}

UHMW-PE grade.

101_{+---r---.---.---.---.----~}

10"1 ₁₀ ₁₀1 ₁₀2 ₁₀3 ₁₀4 ₁₀5

time [s]

Figure 3.8: Stress relaxation modulus (G(t)) of UHMW-PE as a function of time (t) at 180 °C.

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44 Chapter 3

3.4 Discussion

Crystallization from dilute solution, especially if crystallization occurs below the overlap concentration, is an effective way to obtain disentangled polymer systems1_•2_•

The solid state properties of salution-cast LDPE and UHMW-PE, as presented in

figures 3.2 and 3.6, can be easily explained in terms of a strongly reduced entang]ement density. The disentangled situation in salution is preserved in the solid state even after remaval of the solvent as a consequence of crystallization into chain-folded crystals.

If we adopt the classical representation of an entanglement (figure 3.1) no difficulties are encountered in onderstanding the process of disentangling via dissolution. Chains as a whole are separated upon dissalution and in the dilute regime, at polymer concentrations below the critica! concentration for coil overlap, the number of intermolecular cantacts is reduced to the extreme. The surprising observation, however, is the short time scale needed to destray the solution-induced drawing characteristics. Heating above the melting point for a short time renders a solid ( after crystallization) or a melt that are indistinguishable from their equilibrium counterparts (figures 3.2, 3.3, 3.4, 3.6, 3.7). In other words, the time-scale needed for restoring equilibrium properties could not be measured.

The longest relaxation times in both LDPE and UHMW-PE (figures 3.3 and 3.8), are at least two orders of magnitude larger than the observed time-scale of memory-effects in salution-cast samples. This indicates that reptative motion is not a prerequisite to restare equilibrium properties in the solid state and melt. In other words, the experimental data show that segmental motion is sufficient to destray the dissalution history of salution-cast LDPE and UHMW-PE.

Upon melting of solution-crystallized chain-folded PE crystals, the ebains will expand towards a random coil conformation11•12• The driving force will beofan en topic

nature. In order to restare an equilibrium entang]ement density, the formation of entang]ements, as schematically represented in figure 3.1, is virtually impossible since reptative motion of complete ebains is needed at corresponding timescales beyond the experimental times used in this study. The experimental results therefore indicate that intermol ecular interactions, related to drawability and rheology of polyethylenes, cannot be represented by the classica! picture of an entanglement as visualized in figure 3.1.

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Several alternative suggestions have been made with respect to the physical nature of the intermolecular interactions which determine the maximum draw ratio of UHMW-PE. It was proposed that the maximum attainable draw ratio of UHMW-PE in the solid state is related to the degree of adjacent re-entry within crystals13_•_Upon

crystallization from solution, adjacent re-entry is promoted which facilitates unfolding of lamellae during solid state drawing. Heating in the melt, foliowed by a recrystallization results in less adjacent re-entry and consequently a loss in drawability. In this partienlar case segmental motion is sufficient to loose ultra-drawability and consequently a fast decay of ultra-drawability can be explained from a kinetic viewpoint.

Repeatedly extruded, "shear-refined" LDPE samples exhibit long term memory-effects14

-24 which are usually attributed to disentangling of LDPE during

repeated processing. In view of the results concerning solution-crystallized LDPE, these results can nat be simply attributed to restoring an equilibrium entanglement network from a disentangled situation. Shear-refining is less effective, compared to crystallization from dilute solution, to obtain a disentangled system. Moreover, it was already suggested previously by Münstedt18 _{that the phenomena observed in}

shear-refined LDPE are related to alignment of long chain branches along the macromolecular backbone. In subsequent studies, this alternative interpretation of the experimental results on repeatedly extruded LDPE was further substantiated24•

Several routes have been found to produce UHMW-PE grades which are intrinsically ultra-drawable i.e. the so-called "virgin" UHMW-PE grades25•26• The

experimental results in this chapter suggest that the rheological properties in the melt of these virgin UHMW-PE grades are identical to "normal" UHMW-PE grades. Moreover, it is shown that processing of these virgin UHMW-PE grades must be performed in the solid state, in order to preserve their excellent drawability. In fact, solid state routes for the processing of these virgin UHMW-PE grades have been explored27•28• Unfortunately, solid state processing is restricted to fabrication of simple geometries such as (thick) monofilaments, tapes, rods and tubes.

In this study, it was shown that solution-induced properties of UHMW-PE are lost almast instantaneously in the melt. The reverse process, i.e. generating ultra-drawability in melt-crystallized UHMW-PE, bas also been investigated29_-32

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Cbapter 3

found that, upon immersing in a solvent-bath above the dissalution temperature, UHMW-PE films can absorb large amounts of solvent without Ioosing their film geometry. Moreover, in the case of thin melt-crystallized UHMW-PE samples ( <0.1 mm), ultra-drawability could be generated within a few seconds. Basedon a variety of experiments, invalving the influence of molecular weight, swelling temperature, type of solvent and film thickness on the swelling process, it was concluded that the kinetics of the swe1ling process are primarily controlled by solvent diffusion in the films. In other words, diffusion of entire macromolecules through the process of reptation is irrelevant for the re-generation of ultra-drawability in melt-crystallized UHMW-PE films. Therefore, these studies further substantiate the conclusions in this chapter.

The only intermolecular interactions, within the polyethylene crystallattice, are weak Van der Waals forces. Usually, it is assumed that these Van der Waals farces hardly influence the drawability of polyethylenes. In the case of poly(vinylalcohol) comparatively strong hydragen honds exist in the crystallattice. The influence of these hydragen honds on the drawing behaviour and properties of poly(vinylalcohol) fibres is discussed in the next chapter.

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47

3.5 Conclusions

Based on some model experiments, conceming memory-effects in salution cast polyethylenes, it was attempted to identify the physical nature of intermolecular interactions ( chapter 2) in polyethylenes. lt was shown that the dissalution history of linear and long chain branched polyethylenes, with respect to solid state and melt properties, is lost almast instantaneously upon heating into the melt. This indicates that reptation of ebains as a whole is nat a prerequisite for restoring equilibrium properties. Consequently, entanglements, visualized as four strands teading away from a contact, (figure 3.1) are nat appropriate as a representation of the topological constraints which determine the viscoelastic properties of linear and long chain branched polyethylenes in the solid state and melt. An undisputable identification of the intermolecular interactions in polyethylenes was nat possible. However, a suggestion conceming the physical nature of intermolecular interactions was presented.