Oriented structures based on trans-1,4-polybutadiene

Citation for published version (APA):Engelen, Y. M. T. (1991). Oriented structures based on trans-1,4-polybutadiene. Eindhoven: TechnischeUniversiteit Eindhoven. https://doi.org/10.6100/IR362580

DOI:10.6100/IR362580

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

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Oriented Structures

based

on

Trans-l,4-Polybntadiene

PROEFSCHRIFf

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 vcrdedigen op

dinsdag 17 december 1991 am 14.00 uur

door

YVONNE MARIA TIlEODORA TERVOORT-ENGELEN

Geboren te Nijmegen

Dit proef~chrift is goedgekeurd door

de promotoren:

en de co-promotor:

prof. dr. P.J. Lemstra

prof. dr. if. RE.H. Meijer

dr. ing. C.W.M. Bastiaansen

Tout Ie monde est sage apres coup

.Frans spTeekwoord-

Contents

Contents

Chapter 1 General Introduction

1.1 High-Performance Polyethylene Fibres

1.2 The Importance of Intermolecular Interactions for

Fibre Properties

13 Crosslinking of Polethylene Fibres

1.4 Objective of the Thesis

1.4 Survey of the Thesis

1.5 References

Chapter 2 Trans-l,4-Polybutadiene

2.1 Introduction

2.2 Crystal Structures

2.3 Electron Beam Irradiation

2.4 Solid-State Deformation

1

5

9

11

12

13

17

18

21

2.4.1 Uniaxial Tensile Drawing 24

2.4.2 Solid-State Coextrusion 25

2.4.3 Uniaxial Tensile Drawing of Crosslinked T-l,4-PB 26

2.4.4 Deformation Mechanism 28

2.4.5 EB Irradiation of Oriented T-l,4-PB 30

2.5 Conclusions 30

2.6 References 31

ii Contt~nts

Chapter 3 Blending of Trans·l,4·Polybutadiene

and UHMW-Polyethylene in Solution

3.1 Introduction

3.2 Some Aspects of Solution-Crystallization

3.3 Experimental

3.3.1 Solution-Cry~tallization of T-l,4-PB, UHMW-PE and

35

36

their Blends from Xylene 40

3.3.2 Preparation of Blended Films 41

3.3.3 Differential Scanning Calorimetry (DSC) 41

3.3.4 Transmission Electron MicroscoPY (TEM) 42

3.3.5 Wide-Angle X·ray Diffraction (WAXD) 42

3.3.6 Fourier-Transform Infra"Red SpectroscoPY (FT-IR) 42

3.3.7 Quantitative Analysis of the Crystallinity 43

3.3.8 Nuclear Magnetic Resonance (NMR) 43

3.4 Results and Discussion

3.4.1 Crystallization of UHMW·PE from Xylene

3.4.2 Crystallization of T-l,4-PB from Xylene

3.4.3 The Morphology of Blended Films

3.5 Conclusions

3.6 References

Chapter 4 Morphology and Properties of

Drawn UHMW-PE/T-l,4-PB Blends

4.1 Introduction

4.2 Experimental

4.2.1 Sample Preparation

4.2.2 Morphology

4.2.3 Degree of Orientation

44

45

48

57

58

61

62

62

63

Contents

4.2.4 Tensile Testing

4.3 Results and Discussion

4.3.1 Drawability

4.3.2 Morphology

4.3.3 Tensile Properties

4.4 Conclusions

4.5 References

Chapter 5 Electron Beam Irradiation of

Drawn UHMWvPE/T-l,4-PB Blends

5.1 Introduction

5.2 Experimental

5.2.1 Sample Preparation

5.2.2 Electron Beam Irradiation

S.2.3 Gel Fraction and Swelling Ratio

5.2.4 Maximum Draw Ratio

5.2.5 Tensile Testing

5.2.6 Creep Measurements

5.2_7 Differential Scanning Calorimetry

5.2.8 Constrained Heating

5.3 Results and Discussion

5.3.1 Crosslinking Efficiency

5.3.2 Mechanical Properties

5.3.3 Melting Behaviour and Constrained Heating

5.3.4 Residual Properties after Constrained Heating

5.3.5 Transverse Properties

5.4 Conclusions

5.5 References

iii

64

64

65

75

77

77

79

80

81

81

82

83

83

83

84

84

88

94

96

101

102

103

IV

Chapter 6 Network Topology

6.1 Crol:)slinking in Retrospect

6.2 Crosslinking in Prospect

6.3 References

Appendix Morphology of Nascent UHMW-PE:

Chain-Extended vs. Chain-Folded Crystals

A.1 Introduction

A.2 Experimental

A2.1 Materials

A2.2 Differential Scanning Calorimetry

A2.3 Transmission Electron Microscopy

A.2A Average Crystal Sizes

A.2.5 Electron Beam Irradiation

A3 Results and Discllsl:)ion

A.4 Conclusions

A5 References

Summary

Sa menvatting

Nawoord

Curriculum Vitae

Contents

105

108

109

111

112

113

113

113

114

114

118

119

121

125

129

131

General Introduction 1

Chapter 1

General Introduction

1.1 High~Perfol"mance Polyethylene Fibres

Within the present range of commercially available synthetic fibres, solution­

spun/drawn ultra-high molecular weight Polyethylene (UHMW-PE) fibres perform

rather well in view of their impressive tensile properties, viz. moduli in the range

of 100 - 150 OPa and tenacities approaching 4 OPa [1]. Due to the relatively low

density of polyethylene, appro 1000 kgjm3, the specific values for stiffness and

strength of the PE fibres are even more impressive, especially in comparison with

classical fibres based on steel and glass, see figure 1.1 below [2]. Besides a high

(specific) strength and stiffness in the fibre direction, these so-called high­

performance polyethylene (Rr-FE) fibres demonstrate a relatively large work to

break and possess good damping characteristics due to their visco-elas tic nature [1].

In view of this combination of properties, HP-PE fibres are used in ballistic

applications like bullet-proof vests. In structural hybrid composites, HP-PE fibres

are combined with intrinsic brittle fibres, like carbon and glass fibres, to improve

the impact resistance [3,4,5,6].

2 Chapter 1

Currently, Hp·PE fibres are produced by DSM/Toyobo (Oyneema) and its licensee

Allied Signal (Spectra) via solution(gel)-spinning of ultra-high molecular weight

polyethylene. Mitsui adopted the ~ame technique for the production of fibres based

on lower molar mass polyethylenes resulting in inherently less impressive tensile

properties (TekmiJon),

Figure 1.1 Specific strength vs- specific modulus of various fibres

(Nltex '" 1 GPa/ p; p '" d(msity in gJcm3)

(reproduced with pf:rmission from referen(:(~ f2])

The principle of solution(gel)-spinning was found in the late seventies at DSM­

Research [7,8,9,10] and further developed in the eighties into a commercial

spinning process, Figure 1.2 shows schematically the basic features of this process,

focused on the essential ~teps for spinning and drawing. A semi-dilute solution of

ultra-high molecular weight polyethylene, the weight average molar mass typically

above 106 g/mole, is spun into a quench bath, for example water. Upon cooling, PE

crystallizes in folded-chain lamellar crystals which are still :surrounded by solvent

molecules, and a gel-like filam<;:nt is ohtained. Gelation/crystallization is caused hy

General Introduction 3

the fact that the polymer chains in sOlution still contain entanglements (the overlap

concentration for these high molar mass polyethylenes is rather tow, < 1 %), which

are trapped upon solidification and provide cOilllectivity throughout the system.

The filament:s can be drawn in an Oven to high draw ratios, in a temperature range

close to, but below, the dissolution c.q. melting temperature of the system.

Figure 1.2

Continuous e~tlU,lol1/$oll.ltlon. Twln-screw/co-rolsllng

Quenchlng./extractlon btllh oven

suspension UHMW.PE:

fibre

Schematic representation of the solution-spinning process oj UHMW-PE

fibres (reproduced with pennission from reference Ill)

The remarkable feature of thi~ process is, that the as-spun and quenched fibres are

also highly drawable if the solvent is removed (by extraction or evaporation) prior

to the drawing operation [8,11,12,13]_ The effectiveness of the drawing

process, i.e. the slope of the E-modulus vs. draw ratio curve, is not influenced hy

the amount of solvent present in the fibrous system_ A unique relation is found for

the Young's modulus as a function of draw ratio, which is not influenced by

molecular weight, drawing temperature or initial polymer concentration in solution

[8,11,12,13,14, 15,16,17,18].

4 Chapt(~r 1 ----

Apparently, the strongly enhanced drawahility of solution-spun UHMW-PE in

comparison with its melt-crystallized counterpart, can not be attributed solely to a

pla~ticizing effect of solvent present within the as-spun/quenched fibres.

Eqllally important is, that the solvent induces an optimum structure in the as-spun

fibres for ultra-drawing. Once this structure is generated, the solvent can be

removed and drawability is preserved even in the solid state, provided of course that

during removal of solvent no melting or re-dissolution occurs [19l-

The drawing mechanism of solution-spun UHMW-PE fibres was investigated by

Smith et al. [12]. [t was shown that the maximum attainable draw ratio of the as­

spun filaments is inversely proportional to the square root of the initial polymer

concentration in solution. The enhanced drawability of solution-spun/ca."t

polyethylenes was related to the entanglement density in the system. Since chains

can not Cross mutually, entanglements act as friction points during draw. The

dissolution step reduces the entanglement density, about proportional to the relative

amount of solvent. The disentangled state is 'frozen' in the solid state due to

crystallization and the solvent can be removed provided that the crystals remain

intact during this procedure. During drawing at elevated temperatures, in a

temperature range close to but below Till' the polyethylene crystals are rather

ductlle and can be deformed plastically_

Recently, Irvine and Smith demonstrated that the development of the Young's

modulus as a function of the draw ratio could be modelled quantitatively based on

the entanglement-concept. The model is based on the deformation of an entangle­

ment network in which the sub-chains are fully stretched and contribute to the axial

Young's modulus depending on their degree of orientation in the drawing direction

[20]. The crystals provide a viscous matrix which prevents relaxation of the

oriented chains but can be ignored as far as the deformation process is involved.

This approach is adequate in the case of PE in view of the ductile character of the

crystals during drawing in a temperature range: T(> < T~r.w .;; T m'

T(> is the so-called a-transition temperature of polyethylene, ,1 transition which

General Introduction 5 ------------------------~-'" .. ,,---

occurs at appr. 80 ~C and is generally related to the onset of molecular motions

within the crystal~ [21,22,23,24,25,26,27,281-

The high drawability of solution-spun/cast polyethylenes, once entanglements have

been (pardy) removed, is on the one hand an advantage for obtaining chain­

extended structures. On the other hand this ductility imposes severe limitations on

the application of the fibre under prolonged static loading conditions (resulting in

creep) since no lock·in mechanism is operative after the drawing process has been

completed. In principle, there will be no basic difference between drawing (in the

process of making fibres) and creep during axial loadings. Both deformation

processes can be considered to be identical, albeit operative on a different time and

temperature scale [29].

Besides the low creep resistance of PE fibres there are other properties which limit

the range of applications of these fibres) e.g. the relatively low melting temperature

(145 QC) and the low shear moduli. These properties are all governed by the weak

Van der Waals interactions operative between the chains in the oriented structure.

In the following section the importance of intermolecular interactions for fibre

properties will be discussed.

1.2 The Importance of Intermolecular Interactions

for Fibre Properties

The research efforts in the area of oriented polymers, which started in the early

sixties, were stimulated to a large extent by theoretical calculations and estimations

concerning the stiffness and strength of a fully extended polymer chain. It was

calculated that, in the case of a fully extended polyethylene chain, the Young's

modulu5 could amount up to 250·300 OPa [30,31,32,33,34,35,36]

with a corresponding tensile strength, related to rupture of the covaJent C·C bonds

in the main chain, of 20·30 OPa [37,38,39,40,41]. Similar impressive

6 Chapter 1 ------

results were obtained for other "zig-zag" chains like poly(vinyl alcohol) and

polyamides. Polyethylene however seems of particular interest in view of the small

chain diameter due to the absence of pendent groups. If extended polyethylene

chains are packed in a (::rystalline register with the chain axes parallel to the fibre

axis, a high number of load bearing elements i:;; present per fibre cross-section and

consequently impressive tensile properties could be expected.

Although these theoretical calculations and/or estimations concerning strength and

stiffness of polymer chains stimulated research efforts to pursue chain-extension in

polymeric systems, their intrinsic value and predictive power is rather limited. The

calculations are based upon loading infinite chains whereas in practice chains are

finite! The major consequence of a finite chain length is, that the tensile properties

are not only determined by the strength and stiffness of the covalent bonds in the

main chain but also by the secondary bonds, i.e. secondary interactions between

chains_

The importance of intermolecular interactions for the ultimate fibre properties wa."

demonstrated recently by Termonia et aL {41,42,43,44,45]. They describe

the deformation of an array of perfectly oriented and extended polymer chains with

varying molecular weight, molecular weight distribution (M,./Mn) and strength of

the secondary bonds (intermolecular interactions) in the system, see figure 1.3.

Figure 1.3 An array of petfectly aligned and extended polymer chains

General Introduction 7

It appears that at a given deformation rate and temperature, the strength of the

(model) fibres increases with increasing molecular weight, with decreasing

molecular weight distribution, with increasing strength of thc secondary bonds and

with increasing overlap length of the separate chains_

Molecular weight and molecular weight distribution are important because they

determine the number of chain ends in the system. Chains ends act as stress

concentrations, decreasing the tensile strength. The overlap length of the chains,

also related to tbe molecular weight, governs the competition between chain

slippage and chain scission, as failure mechanisms of the fibres. This is illustrated

by figure 1.4 which shows the stress-strain curves of polyethylene fibres as a function

of molecular weight. Clearly a transition from ductile to brittle failure can be

recognized at molecular weights around lOS g/mole [41].

Figure 1.4

12

o 2 3

strain (%)

Stress-strain curves of polyethylene fibres as a junction of molecular

weight (reproduced with pennission from reference /41J)

8 Chapter!

The influence of the strength of the ~econdary bonds is illustrated in figure 1.5

which compares the strength of PPT A and polyethylene fibres as a function of

molecular weight_ In the polyethylene fibre only weak Van der Waals interactions

are operative between the chains whereas in PPTA strong hydrogen bonds are

formed. It appears that to obtain a strength of 5 OPa, for polyethylene fibre!> a

molecular weight of > HY g!mole is necessary whereas the same strength is

obtained for PPTA already at a molecular weight of lit gjmolc.

\LO~:;-· --L----L.--'-J....:..L...-'-'-'..L10--;-4----L.--'-----'---'--'-...L..LLllO--::~,------L-----L--I.. ........ ---'-'--'-'106

molecular lM;!igh1 (g/mole)

Theoretical strength of PPTA and PE fibres as a function of molecular

weight (reproduced with permission from references [41] and f45J)

The parameters mentioned (M,.., M. • ../Mn' strength of secondary bonds and overlap

length) are related to each other. Polymers with strong secondary bonds require a

!;maller overlap length to ohtain a certain degree of intermolecular interactions and

high tenacities can in principle (in case of perfectly aligned systems) be reached at

lower molecLllar weights_

In principle, an increase in intermolecular interactions might also improve the creep

resistance of PE fibre5, as well as the temperature resistance and off-axis properties.

General f ntroduction 9

An obvious way to obtain stronger intermolecular interactions is to start with

macromolecules more polar than PE, like polyamides or poly(vinyl alcohol). Similar

to polyethylene, some of these polymers possess a high theoretical modulus in the

chain direction [46,47,48]. However, a principal disadvantage of these 'polar'

polymers is that due to the presence of strong intermolecular interactions, the

chain-folded crystals are more difficult to deform (slip) and to transform into chain­

extended crystals, even if the chains are largely disentangled. In the case of flexible

polyamides, the hydrogen bonds within the crystalline regions limit drawability to

draw ratios of 4 to 5 [49], whereas in PVOH the maximum attainable draw ratio

is about 20 [50]. And, although at a given draw ratio the tensile properties of

these polymers are higher than of polyethylene, this does not compensate for the

low maximum draw ratio. The net effect of the enhanced intermolecular inter­

actions in PVOH e.g., is a decrease in maximum attainable Young's modulus and

tensile strength « 75 OPa and < 2.5 GPa respectively [50]).

Therefore it seems more favourable to attempt to introduce additional intermo­

lecular interactions in UHMW-PE fibres, e.g., by introducing covalent bonds

between chains via crosslinking.

1.3 Crosslinking of Polyethylene Fibres

As stated above the creep resistance and other less favourable properties of the HP­

PE fibre might be improved by an increase in intermoleCUlar interactions.

This might be realized by the introduction of homogeneously distributed chemical

crosslinks (a homogeneous network) in the fibre. This method is essentially different

from the use of more polar polymers, since in principle the network can be

introduced in already drawn PE fibres without interfering with the drawability of

the system.

In the past several techniques were developed to achieve such networks. One of

these techniques involves chemical crosslinking of PE fibres by using e.g. dicumyl

Chapter J

peroxide (OCP) [51]. However, attempts to introduce DCr into already spun and

drawn PE failed, due to the high degree of orientation and crystallinity of the highly

drawn fibres. Therefore, the crosslinking agent had to be introduced into the as­

spun fibres, prior to drawing, but in this way the drawability and the tensile

properties of the fibres were reduced strongly [51].

Recently, a study about photo-crosslinking of ultra-high strength PE fibres has been

reported [52,53,54]. These fibres possess a high strength (2.5 OPa) but only

a moderate modulus (46 GPa). In this study, the photo-initiating agent benzophe­

none (BP) was introduced in the spun and drawn fibre by vapour absorption at

elevated temperatures (100 QC). Subsequently, the fibres were UV irradiated. At

short loading times, the creep rate of the treated fibres was higher than the creep

rate of the untreated samples but levelled off at prolonged loading times and

became lower than the creep rate of the untreated fibres after 104 min. Though this

route seems promising to improve the creep resistance of the ultra-high strength

fibre studied, it would probably fail for high modulus fibres since in these fibres the

vapour absorption will be very low due to the high degree of crystallinity and

orientation. The crystals as such are impermeahle for crosslinking agents like DCP

and BP.

The most intensively studied method to introduce crOSS links in PE fibres is high­

energy radiation. Both y irradiation and electron beam (EB) irradiation have been

applied in the past. The main disadvantage of the use of high"energy irradiation is,

that besides the formation of crosslinks also main-chain scissioning occurs

[55,56], which reduces the tensile strength of the PE fibres [57,58,59].

The degree and ratio in which main-chain scissioning and crosslinking occur upon

EB irradiation of polymers is determined by the morphology, the chemical structure

and the radiation dose [60]. In the case of gel.~pun PE fibres the high degree of

crystallinity prohibits succes~ via this route, since crosslinking preferentially ()CCUrs

in the amorphous regions [61,62]. The main reason for this phenomenon is, that

within the crystal lattice, the carbon atOmS are too far apart to allow for the

General Introduction 11

formation of primary bonds between adjacent chains.

So far, significant reductions in creep rate by high energy radiation have only been

achieved for melt-~pun fibres, which were irradiated prior to drawing

[63,64,65]. As far as gel-spun UHMW-PE fibres are concerned some improvement

in creep resistance could be obtained by irradiation of the as-spun fibres before

drawing [66] and at intermediate stages of drawing [67]_ However in both cases

this result was only reached at the expense of drawability and tensile properties.

In conclusion, if the ratio crosslinking/main"chain sdssioning, resulting from BB

irradiation of gel-spun PE fibres, could be increased, it should be possible to

introduce a homogeneous network without too much loss of mechanical properties_

This might be realized by adding so-called prorads to the fibres. Prorads are

crosslinking agents containing functional groups which are highly sensitive to EB

irradiation, e.g. acetylene. Woods et al. showed that irradiation of melt.spun fibres

in which acetylene gas was absorbed resulted in an improvement of the creep

resistance without significant damage to the other properties [68,69]. Again this

route was far less successful for gel·spun UHMW-PE fibres because of the

difficulties with penetration of gas into these highly crystalline fibres.

1.4 Objective of the Thesis

In this thesis, oriented structures based on trans-l,4-polybutadiene (t~1,4-PB) are

explored in view of the enhancement of intermolecular interactions in fibrous

systems. Trans-l,4·polybutadiene, a polymer with a molecular structure resembling

linear polyethylene, possesses a high G-value for crosslinking [70} related to an

abundance of unsaturated bonds in the main chain, which can be used to introduce

covalent intermolecular interactions, i.e. chemical crOSs links. Furthermore, t-l,4-PB

has a conformationally disordered (condis) phase between 70 DC and 140 QC, in

which molecular segments are very mobile [71J, which is expected to further

enhance crosslinking via electron beam irradiation.

12 Chapter J

Both the possibilities of oriented structures based un pure t-1A-PB as well as the

possibilities of oriented products based on blends of t" 1,4.PB and UHMW-PE were

investigated.

1.5 SOn'ey of the Thesis

The contents of the following chapters are briefly summarized below.

In chapter 2 the deformation behaviour and properties of both melt- and solution·

crystallized trans-l,4-polybutadiene films will be presented. Furthermore, cross­

linking of the oriented structures will be discussed.

Chapter 3 describes the solution.blending of t-l,4-PB and UHMW-PE and the

resulting morphology of the solution-crystallized blends as a function of t.l,4-PB

content.

Drawing of the solution-crystallized blend'\, and the resulting fibre properties, in

relation to the fibre morphology, will he discussed in chapter 4_

In chapter S the network formation by EB irradiation in drawn UHMW.PE/t-l,4-

PB blends is discussed.

Finally, in chapter 6 the results of the previous chapters are interpreted in view of

the objective of the thesis,

In the appendix some aspects of the morphology of nascent UHMW-PE reactor

powders will be discussed, focusing on the question whether chain.extended or

chain-folded crystals prevail in these samples,

General Introduction 13

Crosslinked oriented strnctures with improved thermal resistance, based On UHMW­

PE/t-l,4-PB blends have been patented by DSM [72}.

The appendix to this thesis has recently been published [73 J and the author has

contributed to some papers on subjects not presented in this thesis

[74,75,76, 77J.

1.6 References

1. l..emstra, PJ., Kirschbaum, R., Ohta, T. and Yasuda, H., 'Developments in

Oriented PolymeT3-2" Ed. I.M. Ward, Elsevier Appl. Sci. Publ., New York,

1987, chapter 2

2. Lemstra, P J. j van AerIe. NAI.M. and Bastiaansen, C.W.M., Polym.l 1987.

19,85

3. Peijs. A.AI.M. and Lemstra, P J., 'Integration of Fundamental Polymer

Science and Technology', pan 3, Ed. P.I. Lemstra and LA Kleintjens,

Elsevier Appl. Sci. Publ., London, 1989. p. 218

4. Peijs, AAJ.M., Catsman, P., Govaert; LE. and Lemstra, P J., Composites

1990,21, 513

5. Peijs, AAJ.M., Venderbosch, RW. and Lemstra, P 1., Composites 1990,21,

522

6. Peijs, A.A.J.M. and Venderbosch, RW.,J. Mater. Sci. Lett. 1991, 10, 1122

7. Smith, P. and Lemstra, P.I., GB Patent 2.051.667, 1980

8. Smith, P., Lemstra, P.I., Kalb, B. and Pennings, A.J., Polym. Bull. 1979, 1,

733

9. Lemstra, P.J. and Kirschbaum, R, Polymer 1985, 26, 1372

10. Lemstra, PJ., Bastiaansen, C.W.M. and Meij<;:r, H.E.H., Die Angew.

Makromol. Chern. 1986, 145/146, 343

11. Smith, P. and Lemstra, P.1., Polymer 1980, 21, 1341

14 Chapter 1

12. Smith, P., l~mstra, P_I_ and Booij, H.C., 1. Polym, Sci. Polym. Phy.,. Ed. 1981,

19,877

13. Smith, P_ and Lemstra, P.I., Coil. Polym. Sci, 1980,258, 891

14. Capaccio, G. and Ward, I.M., Nature (phys, Ed,) 1973,243, 143

15. Capaccio, G. and Ward, I.M., Polymer 1974, 15,233

16. Capaccio, G. and Ward, I.M., Polymer 1975, 16, 239

17. Capaccio, G. and Ward, I.M" Polym. Eng. Sci. 1975,15,219

18. Capaccio, G_, Crompton, T.A. and Ward, I.M.,J. Polym- Sci. Polym. Phys. Ed.

1976, 14, 1641

19. LemstTa, P,J_, Bastiaansen, C.W.M. and Meijer, H.E.H., Angew. MakromoL

Chern. 1986, 145/146,343

20_ Irvine, PA and Smith, P., Macromolecules 1986, 19, 240

21. Mansfield, M. and Boyd, R.H.,J. Polym. Sci- Polym. Phys. Ed. 1978,16, 1227

22. Olf, H.O. and Peterlin, A, 1. Polyrn. Sci. A-2 1970,8, 753

23. Olf, H.G. and Peterlin, A,l. Pofym. Sci. A-2 1970,8, 771

24, Opella, S-J- and Waugh, J.S., 1. Chern. Phys- 1977,66,4919

25. Reneker, D_H. and Mazur, J., Polymer 1982,23,401

26. Ewen, B., Strobl, T.R. and Richter, D., Faraday Disc. Chern. Soc. 1980, 69,

19

27. Boyd, R.H., Polymer 1985, 26,323

28. Boyd, R.H., Polymer 1985, 26, 1123

29. Oovaert, LE., Ph.D. Thesis, Eindhoven University of Technology, The

Netherlands, 1990, chapter 4

30. Shimanouchi, T_, A~ahina, M. and Enomoto, S_,l. Polym. Sci. 1962, 59, 93

31. Odajima, A. and Maeda, T., J- Polym. Sci. 1966, IS, 55

32. Wobser, G. and Blasenberg, S., Coil. Polym, Sci. 1970,241, 985

33. Perepelkin, K,E., Di(~ Angew. Makromol. Chem- 197"1,22, 181

34. Broudeaux, D.S.,J Polym. Sc.:i. Polym. Phys. Ed. 1973, ll, 1285

35. Tashiro, K., Kobayashi, M. and Tadokoro, n., Macromolecules 1975, 11,914

General Introduction 15

36. Christ, B., Rather, M.A., Brouwer, AL. and Sabin, J.R.,1 Appl. Phys. 1979,

50,6047

37. Smith, P. and Lemstra, P.J.,1. Polym. Sci. Polym. Phys. Ed. 1980, 19, 1007

38. Smith, P. and l..emstr~ P J., J. Polym. Set. Pofym. Phys. Ed. 1982, 20, 2229

39. Smook, J., Hamersma, W. and Pennings, AJ., 1. Maler. ScL 1984, 19, 1359

40. Wagner, RD. and Steenbakker, L W., PhiL Mag. Lett. 1989, 59, 77

41. Termonia, y, Meakin, P. and Smith, P., Macromolecules 1985, 18,2246

42. Termonia, Y. and Smith, P., Macromolecules 1987,20,835

43. Smith, P. and Termonia, Y., Polym. Comm. 1989,30,66

44. Termonia. Y. and Smith, P., Macromolecules 1988, 21, 2184

45. Termonia, Y. and Smith, P., Polymer 1986, 27, 1845

46. Perepelkin, K.E., Die Angew. Makromol. Chern. 1971,22, 181

47. Tashiro, K.., Kobayashi, M. and Tadokoro, H., Macromolecules 1978, 11,914

48. Tashiro, K, Tadokoro, H., Macromolecules 1981, 19,481

49. Postema, AR., Smith, P. and English, AD., Polym. Comm. 1990, 31, 444

50. Bastiaansen, C.W.M., Ph. D. Thesis, Eindhoven UniverSity of Technology,

The Netherlands, 1991

51. de Boer, J., Van den Berg, H.J. and Pennings, AJ., Polymer 1984, 25, 513

52. Chen, Y.L. and R~.nby, B.,i. Polym. Sci. Part A: Potym. Chern. 1989,27,4051

53. Chen, Y.L and Rinby, B., 1. Polym. Sci. Part A: Polym. Chern. 1989,27,4077

54. Chen, Y.L and Rllilby, B., Polym. Adv. Techn. 1990, 1, 103

55. Charlesby, A, 'Atomic Radiation. and Polymers', Pergamon Press, Oxford,

1960

56. Dole, M., 'The Radiation Chemisrryoj Macromolecules', Academic Press, New

York, 1972

57. de Boer, J. and Pennings, A.J., P{)lym. Bull. 1981, 5, 317

58. de Boer, 1. and Pennings, A.J., Coli. Polym. Sci 1983,261, 750

59. Dijkstra, 0.1. and Pennings, AJ., Polym. Bull. 1987, 17,507

60. Keller, A., 'Developments in Crystalline Polymers-]', Ed. D.C. Bassett, AppL

16 Chapter 1

Sci. Publ., London, 1982

6L Patel, G.N. and Keller, A., 1. Polym. Sci. Polym. Phys. Ed. 1975, 13,303

62. Jenkins, R and Keller, A, J. Macromol. Sci. Phys. 1975, B11, 301

63. Wilding, M.A. and Ward, LM., Polymer 1978, 19,969

64. Ward, LM. and Wilding, M.A.,l. Polym. Sci. Polym. Phys. Ed. 1984,22, 561

65. Klein, P.G., Ladizesky, N.H. and Ward, I.M., J. Polym. Sci. Polym. Phys. Ed.

1986, 24, 1093

66. Hikmet, R, umstra, P.l. and Keller, A, ColI. Polym. ScL 1987, 265, 185

67. van AerIe, N.AJ.M., Crevecoeur, G. and Lemstra., P.J., Polym. Comm. 1988,

29, 128

68. Woods, D.W., Busfield, W.K. and Ward, I.M., Polym. Comm. 1984,25,298

69. Woods, D.W., Busfie1d, WK and Ward, I.M., Plast. Rubb. Process. Appl.

1985,5, 157

70. van Gisbergen, J.G.M., Ph.D. Thesis, Eindhoven University of Technology,

The Netherlands, 1991, chapter 2

71. Moller, M .• Makromol. Chem. Rapid Comm. 1988,9, 107

72. DSM Stamicarbon, NL 9.001.069, 1990

73. Tervoort-Engelen, Y.M.T. and Lemstra, P.I., Polym. Comm. 1991,32,343

74. Tervoort-Engelen, Y.M.T. and van Gisbergen, J.G.M., Potym. Comm. 1991,

32,261

75. Vandeweerdt, P., Derghmans, H., Tervoort. Y., Macromolecules 1991, 24,

3547

76. Gerrits, N.SJ.A., Tervoort, Y., accepted for publication in 1 Mater. Sci.

77. Heynderickx, L, Broer, D.l. and Tervoort"Engelen, Y., accepted for

publication in .I. Mater. Sci.

Trans-I ,4-Polybutadiene 17

Chapter 2

Trans-l,4 .. Polybutadiene

2.1 Introduction

Polybutadiene is an unsaturated polymer with four isomers, see figure 2.1 [1]. For

the production of strong and stiff fibres, trans-1,4-polybutadiene (t-l j 4-PB) is the

most promising candidate amongst these four. It is a linear polymer with no

pendent side groupS and a comparatively high melting temperature (145 cC), rather

similar to the melting temperature of polyethylene.

Traditionally t-l,4-PB is synthesized via Ziegler-Natta polymerization of 1,3-

butadiene. The t-l,4-PB used in this thesis is polymerized in our laboratory, at room

temperature using a VCl3/ Al (C2HSh catalyst system [2].

In this chapter some possibilities for the production of oriented t-l,4-PB structures

are introduced. Literature concerning the crystal structure, the effects of electron

beam irradiation and solid"state deformation of t-l,4·PB will be reviewed.

Furthermore, some new results will be presented concerning uniaxial drawing of t-

1,4-PB followed by EB irradiation.

18 Chapter 2

( -··CH CH -\ 2'-. ./ 2 'n

CH=CH cis-l,4-PB (2 '1:::)

(-CH2'-...

CH~CH'-... trans·1,4-PB (145°C)

CH2-)n

(-CH -·_·CH 2 I

CH II

CH 2

lr 2

CH syndiotactic 1,2-PB (126 Oc) I

('·~CH2-CH-CH2-CH- CH2-CH-----7n I I

CM CH II II

Figure 2.1 Isomers of polybutadiene with their respective melting temperatures

2.2 Crystal Structures

At room temperature t-1,4-PB possesses a monoclinic crystal structure. The unit cell

with lattice constants a:::;; 8.63 A. b = 9.11 A. and c = 4.83 A, includes four molecular

segments. This is larger than the orthorhombic unit cell of linear polyethylene

(a=7.40 A, b =4_94 A. and c=2.54 A) which also includes 4 molecular segments_ The

chain conformation of t-l,4-PB is not a. planar zig-zag as encountered in poly­

ethylene, since the bonds adjacent to the double bonds are rotated out of the trans

conformation hy + 7r and - 71°, leaving the central CH2-CH2 group in the low"

energy trans conformation [3] (see figure 2.2).

Trans·114·Polybutadiene 19

Figure 2.2 Chain confonnation oftrans-l,4-po!ybutadiene in the monoclinic lattice

Upon heating t-1,4-PB, a first-order solid-solid phase transition occurs from the

monoclinic low-temperature crystal structure into a hexagonal crystal structure. This

transition was reported first in 1956 by Natta et al. [4].

The transition temperature is dependent on crystallization conditions [5,6,7] and

pressure [8,9]. Values ranging from 48 "C to 76 "C have been reported

[5,6,7,8,9,10,11,12,13,14].

The hexagonal phase melts at about 140 to 145°C which is comparable to the

melting temperature of polyethylene.

Both transitions, the solid~solid phase transition and the melting peak, can be

observed by differential scanning calorimetry (OSC) as is illustrated in figure 2.3.

The theoretical heat of fusion and the heat of transition amount to 69 Jig and

144 Jig respectively (12).

20 Chapter 2

140 10

temperaHlre eC}

Figure 2.3 DSC thermogram of trans"1,4·polybutadiene crystallized from toluene.

The solid-solid phase transition involves an increase in chain separation of about

7.5 % and a decrease in chain length of 4 %, resulting in an increase in overall

volume of + 9 %. The relative chain position and orientation however, do not

change [15,16,171.

The high-temperature crystal form is a so-called conformationally disordered

(condis) phase. In this phase, the posi tional and orientational order of the molecules

are largely preserved, as in a crystal, but each repeating unit can adopt more than

one conformation. M a r:esult the crystal is partially disordered in the chain

direction and some uni·directionalliquid-like motion is possible. The relatively high

molecular mobility of t-l,4-PB in the hexagonal phase was demonstrated yja nuclear

magnetic resonance by Jwayanagi and Miura in the sixties [18] and recently

studied in detail by Moller [19,20]. For more details concerning the candis

phase in general the reader i~ referred to comprehensive reviews on this subject

[21,22].

Trans-J,4-Polybutadiene 21

The crystal modulus of t-1,4-PB at rOOm temperature was determined by Nakamae

[23] from X-ray experiments on solid-state coextruded single crystal mats with an

extrusion draw ratio (EDR [24,25]) of about 20. The modulus of the monoclinic

crystal is estimated to be 90 GPa, for the hexagonal crystal it is not known. The

value of 90 GPa for the monoclinic crystal modulus is rather low compared with the

crystal modulus of 235 GPa reported by Nakamae [26] for polyethylene. However,

the lower E-modulus can be understood in terms of the non-planarity of the t-1,4-

PB chain confOIrnation and the higher cross-sectional area per chain compared with

polyethylene.

2.3 Electron Beam Irradiation

Unsaturated polymers in general possess a stmng tendency towards crosslinking,

which is demonstrated by the overall G~value for crosslinking (G(x» of poly­

butadienes which is in the range of 1.55 to 5.8 [27]. The G-value represents the

number of changes, e.g. crosslinks (G(x», or chain scissions (G(s), realised per 100

eVabsorbed energy.

Parkinson and Sears [28] studied the effects of electron beam irradiation on the

olefinic groups in four types of polybutadienes, one of which was a 95 % trans" 1,4-

polybutadiene. From infra-red measurements prior to and after irradiation at room

temperature, they concluded that the most pronounced effect of irradiation in a.Il

types of polybutadienes, is a reduction in the concentration of olefinic groups. In the

(crystalline) high-trans polymers an additional effect is the destruction of

crystallinity.

Although the destruction of the trans olefinic groups in the high.trans polybutadieoe

possesses overall G·values between 11 and 22, the nature of the product formed

through the high consumption of unsaturated groups is not completely known. The

yield of hydrogen is low (G=OA to 0.5) and crosslinking shows yields of only G",,2

to 6. Further possible reactions are the formation of cyclic structures and of

conjugated unsaturated groups [28J.

22 Ch(.lpter 2 ----------------------

The presence of side vinylene-groups can also play a role in the crosslinking

behaviour of polybutadienes since the G-value for the reaction of these groups

amounts to 40 [28]. However, since the concentration of -CH2 =CH2 side.groups in

the t-l,4-PB used in our studies is only about 1% (according to BC NMR) and the

reaction rate is first order dependent on the concentration, this contribution is not

expected to be a dominant factor in the irradiation of t·l,4.PB.

To the author's knowledge, irradiation of t-l,4-PB in the hexagonal phase has not

been studied explicitly. COnsidering the high segmental mobility of t-l,4.PB chains

in the hexagonal phase, a high crosslinking efficiency is expected, which could

promote crosslinking within the condis phase. This expectation is based on results

reported by Ungar and Keller concerning EB irradiation of paraffins like

n-tricosane (C2]H4II) [291. They observed that irradiation of these compounds at

temperatures at which the paraffins are in their so·called 'rotator' phase

(comparable to the condis phase), results in drastic increases in crosslinking

efficiency compared with irradiation at lower temperatures. Even irradiation in the

melt was less effective.

In figure 2.4 some data concerning the soluble fraction (sol.fraction) of EB

irradiated t·l,4.PB films as a function of irradiation dose, are plotted in a so-called

Charlesby-Pinner plot [301. The data are derived from gel fraction measurements

on t-l,4-PB samples irradiated at 100 °C, reported in reference [31}. According

to Charles by and Pinner [30], the !)ol·fraction (5) is related to the irradiation dose,

R (in MRad), via equation 2. L

G( ) l.602·1O-18N A S + S l/2 = S + --=-c--,----,-~~

2G(x) G(x)M",·R (2.1)

where NA is the Avogadro constant. A plot of 5 + 51/2 against llr should yield a

straight line. From the slope and the intercept of this line, G(s) and G(x) can be

calculated.

I :t

(/)

+ (/)

2.00 ~--------~-~-------,

1.50

1.00

0.50

I I

0.00 L-__ ----L-__ -----l_~_....l...._~ _ ____L __ ~

0.00 0.10 0.20 0.30 0.40 0.50

23

Figure 2.4 Charlesby-Pinner plot for EB irradiated t-l,4-PB films; (0) melt­

crystallized and (_) solution-crystallized

Figure 2.4 reveals that the data concerning melt-crystallized t-l,4-PB samples yield

a straight line, in contrast to the data concerning the solution-crystallized samples_

A similar deviation from the Charlesby·Pinner plot bas been observed before, e.g.,

for irradiated UHMW-PE fibres [32]. It can be caused by inhomogeneous cross­

linking of the solution-crystallized samples, due to the high degree of crystallinity.

For the melt-crystallized samples a G(x) value of 4.2 and a G(s) value of 0.9 can

be calculated (~;;:: l.3-Hf gjrnol). This implies a high crosslinking efficiency of

t-l,4·PB in the hexagonal phase.

24 Chapter 2

2.4 Solid-State Deformation

2.4.1 Uniaxial Tensile Drawing

Orientation of t-l,4-PB via uniaxial tensile drawing has been shown to be not very

successful. In the past, several authors tried to achieve oriented t-l,4-PB structures

for crystallographic studies. lwayanagi [3] obtained a maximum draw ratio of 5 via

tensile drawing of solution-crystallized t-l,4-PB films at 40 "C, Le. in the monoclinic

cryr.tal phar.e. According to Natta [33] tensile drawing in the hexagonal phase or

in the melt results in draw ratios of only 2.

Recently, van Aerle et al. [34,35] investigated the uniaxial tensile drawing of

t-l,4-PB in more detaiL Their results are reproduced in figure 2.5, in which the

maximum attainable draw ratio of melt-crystallized t-l,4-PB is plotted as a function

of the tensile drawing temperature.

I Q ..-~

~ m -e -;

~

Figure 2.5

25 ~---------------------------------.

20

15

10

5 cr- ..0.--1]

(T ~ -G- - -9 - - - - - EJ

1 25 50 75 100 125

drawing temperature tOe)

The maximum attain.able draw ratio oj melt-crystallized t-],4.PB films

as a function of tellsile drawing temperature (reproduced with permission

from reference /35/)

Trans-J,4-Polybutadiene 25

These results are in agreement with the earlier observations of Iwayanagi [3] and

Natta [33] on solution-crystallized samples. In both cases, melt- and solution­

crystallized samples, high draw ratios could not be obtained due to premature

sample failure.

Under special conditions linear polyethylene also exhibits a conformationally

disordered hexagonal phase which resembles the hexagonal phase of t-l,4-PB

[21,22]. For the tensile drawing of linear polyethylene in this hexagonal phase

similar observations have been reported as for t-l,4-PB. UItra-dravm linear

polyethylene fibres can not sustain any applied stress in the hexagonal condis phase

due to an increased molecular mobility [36,37,38,39]. The chains in the

hexagonal phase do not withstand the applied stress and chain slippage occurs.

2.4.2 Solld"State Coextrusion

A possible way to prevent premature sample failure during deformation is solid­

state coextrusion [24,25] instead of tensile drawing. In this technique, a film of the

materiaJ, placed in between a (split) billet of another polymer, is forced through a

corneal die. This technique has also been applied successfully to poly-ethylene

reactor powders to obtain oriented structures [40,41,42]. The mOJphology of

these reactor powders has been subject of extensive discussions and will be

described in an appendix to this thesis.

Via solid-state coextrusion, van Aerie et aI. [34,35] deformed t-l,4-PB single crystal

mats in the hexagonal condis phase to high extrusion draw ratios (see figure 2.6).

From elastic recovery data it appeared that the macroscopic and molecular draw

ratios are virtually identical, which indicates that coextrusion in the hexagonal

crystal phase is an effective method to obtain chain orientation and extension in

t-l,4-PB. Furthermore, these results imply that, despite the high molecular mobility

in the hexagonal phase, no significant relaxation of orientation occurs during or

after coextrusion of t-l,4-PB.

26 Chapter 2 ----------_ ... _ ..... -.... _------------

The maximum extru~i(ln draw ratio obtained via coextru&ion i& about 23, which

results in a modulus of about 20 GPa and a tensile strength of about 0.4 GPa-

1-.Q ..., ~

~ -b X ro E

Figure 2-6

25

-'" --'(] ,/D

20 ,-,/

,/ /

,/

ral

15 I J

I r;;{ J

10 I I

I

6 5 o---G-'[]

1 25 50 75 100 125

extrusion temperatue ("C)

Maximum attainable extrusion draw ratios of t-l,4-PB single crystal mats

as a junction of solid-state coextrusion temperature (reproduced with

pennission from reference [35 J)

2.4.3 Uniaxial Tensile Drawing of Crosslinked T-l~4-PB

The lack of intermolecular interactions, prohibiting stress transfer between the

chains, seemS the cause of the premature sample failure observed during tensile

drawing of t-1,4-PB in the hexagonal phase.

In order to introduce more coherence between the chains, an attempt has been

made to cro~~link t-l,4-PB films, before drawing, by electron beam (EB) irradiation

[31]. The irradiated t-l,4·PB samples show an enhanced drawi.ng behaviour

compared with unirradiated samples (see figure 2.5). The drawability of both

Trans-l,4-Polybutadiene 27

solution-crystallized and melt-crystallized t-l,4-PB films as a function of irradiation

dose is given in figure 2.7. The maximum obtainable draw ratio via this route is the

same as can be obtained via solid-state coextrusion (see figure 2.6).

o .~

~

Figure 2-7

30 .-------------------------------------,

20

10

1

It \

o 20

\ '. \

40 60 80 100 120

irradiation dose (kGyl

The maximum attainable draw ratio via tensile drawing at 100°C oj EB

irradiated t-l,4.PB films; (.) solution-crystallized and (D) melt­

crystallized

A relatively low irradiation dose (20 kGy) is sufficient to make uniaxial tensile

drawing possible_ From gel-fraction determinations it appears that the t-l,4-PB

samples can be deformed via tensile drawing as soon as a network is introduced

(irradiation dose > gel-point dose), and at this gel-point dose the drawability is

maximum_ Higher doses yield lower drawabilities (see figure 2,7),

The Young's modulus of the oriented t-l,4-PB tapes increases with draw ratio as is

illustrated in figure 2.8. The maximum attainable modulus is 18 GPa at a draw ratio

of 20, the maximum tensile strength is 0.7 GPa.

28 Chapter 2 .,.'''"/"'1''·'''1.",''''·'·'

20

--II ,. .--

" ro 15 .--

~ ,.

.--.--(fJ .... ,. ::J • .--

:a .-,-

0 10 .-E ,.

,-,. rJ)

-- • "~ ,. .'11

>- 5 ,.. ,.

--,.

• 0

5 10 15 20 25

draw ratio (-)

Figure 2.8 Young's modulus of t"i,4-PB tapes as a junction of draw ratio

2.4.4 Deformation Mechanism

The deformation mechanism of H,4-PB has heen investigated on solid-state

coextruded tapes [34,35] and on tensile drawn irradiated tapes [31] using X-ray

scattering techniques and 13C NMR.

From the X-ray experiments it appears that in the initial stages of drawing the

chain-axes rotate from perpendicular to the film surface, towards parallel to the

drawing direction (rotation of about 90 QC). Subsequently, the degree of orientation

increases with increasing deformation ratio, which appears from the gradual

contraction of the reflection arcs in X-ray diffraction patterns. The W AXD patterns

also show that the t-l,4-PB tapes return to the monoclinic crystal structure after

cocxtrusion in the hexagonal phase.

Trans" 1,4-Polybutadiene 29

In figure 2.9 Be NMR spectra of oriented t-l,4-PB tapes are given as a function of

deformation.

Figure. 2.9

a b c

40 30 ppm

l3e NMR spectra of roller-drawn t-l,4.PB tapes as a junction of

de/onna/ion ratio

Comparison of spectra a to d reveals that with increasing deformation, the fraction

of material present in the monoclinic phase increases at the expense of the fraction

of amorphous and/or hexagonal material This implies that the ordering of the

chains increases with increasing deformation ratio [31].

30 Chaptc~r 2 ---_ ... _------- ------_._ ... _--

2.4.5 ER Irradiation of Oriented T-t,4-PB

In an attempt to improve the mechanical properties of the oriented t-I,4-PH

samples by introducing more intermolecular interactions, some drawn samples were

irradiated at 100 T. Samples possessing an initial modulus and tensile strength of

18 OPa and 0.7 GPa respectively, show a modulus and tenacity of 14 GPa and 0.5

GPa respectively, after irradiation (50 kGy, 100 QC). No improvement of the tensile

properties could be achieved_

The slight decrease in tensile properties is probably due to relaxation effects during

the irradiation at 100 ~c.

2.5 Conclusions

Uniaxial drawing of t-l,4-PB in its condis phase is not possible as such, because the

chains in the hexagonal phase can not withstand the applied tensile stress and chain

slippage occurs resulting in premature sample failure.

However, oriented structures of t-l,4-PB can be obtained by solid-state coextrusion

of solution-crystallized films and by tensile drawing of EB irradiated t-l,4-PB films,

in which a loose network has been introduced.

The final properties of the oriented products obtained via these techniques are a

maximum Young's modulus and tenacity of about 20 GPa and 0.7 GPa respectively_

Post-irradiation of the oriented t-l,4-PB tapes does not improve these properties.

The ultimate tensile properties of t-l,4-PB are limited due to several factors. The

chain conformation is non-planar, reducing the crystal modulus, appro 90 GPa

according to Nakamae [23]. Moreover, it is rather difficult to synthesize linear and

high molecular weight t-l,4-PB samples. The latter factor limits both the modulus

and the tensile stength_

Nevertheless, t-l,4-PB is an interesting polymer in view of the abundance of

unsaturated hands in the main chain and the presence of a condi:s phase at elevated

Trans-i,4-Polybutadiene 31

temperatures. Therefore, t"1,4·PB could be used as a macromolecular prorad in

blends with UHMW-PE. In the next chapters this possibility will be discussed.

2.6 References

1. 'Encyclopedia of Polymer Science and Engineering', 2nd Ed., Wiley

lntersdence, New York, 1986

2. Natta, G., Porri, L, Corradini, P. and Morera, D., Chim. e. Ind. 1958,40,

362

3. Iwayanagi, S., Sakurai, L, Sakurai, T. and Seto, T., 1. MactomoL Sci. 1968,

Bl,163

4. Natta, G., Corradini, P. and Porri, D., Rend. Accad. Nazi. Lined 1956, 20,

728

5. Bermudez, S.F. and Fatou, J.M., Bur. Pol. J. 1972,8,575

6. Marchetti, A and Martuscelli, Eo, J. Polym. Sci. Polym. Phys. Ed. 1976, 14,

323

7. Tseng, S. and Woodward, AE., Macromolecules 1982, 15,343

8. Kijima, T., Imamura, M. and Kusumoto, N., Polymer 1976, 17,249

9. Bautz, G., Leute, U., Dollhopf, W. and Hagde, P.C., Call Polym. Sci.

1981, 259, 714

10. Porri, L, Corradini, P. and Morero, D., Chim. e. Industr. 1958,40,362

11. Baccaredda, M. and Butta, E.,1. Polym. Sci. Polym. Lett. Ed. 1961,51,539

12. Finter, J. and Wegner, G., MakromoL Chern. 1981, 182, 1859

13. Grebowicz, J., Aycock, W. and Wunderlich, B., Polymer 1986,27,525

14. Danusso, F., Polymer 1967, 8, 281

15. Natta, G. and Corradini, P.,l. Polym. Sci. 1959, 39, 29

16. Corradini, P., J Polym. Sci. Polym. Lett. Ed. 1969, 7, 211

17. Suehiro, K. and Tagayanagi, M.,l. Macromol. Sci. Phys. 1970, 84,39

32 Chapter 2

18. Iwayanagi, S. and Miura, J., Repts. Progr. Polym. Phys. Japan 1965,8,303

19. Moller, M., ACS Polym. Prepr. 1987,28,395

20. Moller, M., MakromoL Chern. Rapid Comm. 1988,9, 107

21. Wunderlich, B. and Grebowicz, J., Adv. P{)~ym. Sci. 1984, 60/61, 1

22. Wunderlich, B., Moller, M., Grebowicz, J. and Baur, H., Adv. Polym. Sd.

1988,87, 1

23. Nakamae, K., University of Kobe, Japan, pn'vate communication

24. Griswold, p.o., Zachariades, AE. and Porter, R.S., Polym. Eng. Sci. 1978,

18, 1978

25. Zachariades, A.B., Griswold, P.O. and Porter, RS., Polym. Eng. Sci. 1979,

19, 441

26. Nakamae, K. and Nishino, T., 'Integration of Fundamental Polymer Science

and Technology', part 5, Ed. P.1.Lemstra and LA Kleintjens, Elsevier

Appl. Sci. Publ., London, 1990

27. van Gisbergen, J.G.M., PhD Thesis, Eindhoven University of Technology,

The Netherlands, 1991, chapter 2

28. Parkinson. W.W. and Sears, W.c., 'Irradiation of Polymers', Ed. RF.

Gould, Am. Chern. Soc. Publications, U.S.A., 1967, chapter 5

29. Ungar, G. and Keller, A, Polymer 1980, 21, 1273

30. Charlesby, A. and Pinner, S.H., Proc. Roy. Soc. 1959, A249, 367

31. Vossen, R.H.R, M.Sc. thesis, Eindhoven University of Technology, The

Netherlands, 1990

32. Dijkstra, O.J. and Pennings AJ., Po/ym. Bul/. 1987, 17,507

33. Natta, G., Pegoraro, M. and Cremonesi, P., Chim. e In.dustr. 1965,47, 722

34. van Aerie, N.A.J.M., Lemstra, PJ., Kanamoto, T. and Bastiaansen,

C.W.M., Polymer 1991, 32, 34

35. van AerIe, PhD. Thesis, Eindhoven University of Technology, The

Netherlands, 1989, Chapter 7

36. Lemstra, P.J., Bastiaan:Sl;:n, C.W.M. and Meijer, RE.H., Angew.

Makromol. Chern. 1986, 145/146,343

Trans-i,4-Polybutadiene

37. Lemstra, P.J., van Aerie, N.AJ.M. and Bastiaansen, C.W.M., Polym. 1.

1987, 19,85

33

38. van AerIe, N.AJ.M. and Lemstra, P.l., Makromol. Chem. 1988, 189, 1253

39. Bastiaansen, C.W.M. and Lemstra, P.J., MakromoL Chern. MakrornoL

Symp. 1989, 28, 73

40. Kanamoto, T., Tsuruta., A., Tanaka, K., Takeda, M. and Porter, RS.,

Polym.l 1983, 15,327; Macromolecules 1988, 21,470

41. Kanamoto, T., 'Integration of Fwzdamental Polymer Science and

Technology', pan 5, Ed. P.l. Lemstra and LA. Kleintjens, Elseviers Appl.

Sci. Pub]., London, 1990

42. Smith, P., Chanzy, H.D. and Rotzinger, B.P., Polym. Comm. 1985,26,258

Blending of Trans-l,4-Polybutadiene and UHMW-Polyethylene in Solution 35 .

Chapter 3*

Blending of Trans .. l,4·Polybutadiene

and UHMW-Polyethylene in Solution

3.1 Introduction

Blending of two or more polymers is a widely applied method to obtain a material

which exhibits new or improved properties. The resulting properties are largely

determined by the degree of miscibility/compatibility of the constituents and the

related morpbology_ True miscibility of polymers in a blend is rare and in general,

phase separation occurs on a microscopic scale. Blends in which at least one of the

components is crystallizable, form a unique class of polymer blends. The

compatibility question in such blends is rather complex due to the existence of

several phases. Two crystallizable polymers, when blended, Can form mixed crystals

(co-crystallization) or separate crystals embedded in a compatible or incompatible

amorphous phase. The ultimate morphology depends largely upon the crystallization

conditions, e.g., crystallization temperature, composition and, in the case of solution"

* Part of this work was perjonned in cooperation with H. Declan.ann, Inst- for

Makromol. Chemie, Albert-Ludwigs-Univ., Freiburg and with Prof M Moller,

Dept. of Chern. Techn., Twente Univ., Enschede

36 Chapter -'

crystallization, the concentration. Another important factor governing the

morphology of a crystallizable blend is the miscibility of the polymer constituent:-;

in the melt or in solution. Liquid-liquid (L-L) phase separation, prior to

crystallization (L-S), will have a major influence on the final morphology.

In the case of mixing trans-l,4-polybutadiene and UHMW-polyethylene with the

aim of preparing blended fibres via the gelspinning technique, mixing must take

place in solution in order to obtain largely disentangled, (ultra)drawablc precursors.

Since the major part of this thesis describes studies performed on blended tapes

obtained via drawing of solution-crystalJj~ed films, the crystallization of both

componems in xylene and the morphology of the dried blended films was

investigated. The crystallization conditions used were preset by the preparation

method of the films.

3.2 Some Aspects of Solution-Crystallization

Crystallization of polymers from solution can roughly be divided in two extremes,

i.e., crystallization from dilute solutions and crystallization from concentrated

solutions.

Upon cooling (semi-)dilute solutions (where the concentration, (/), is higher than the

overlap concentration, ¢ ,), crystallization occurs at a certain degree of supercooling

.6. T ( = Tin . Tc) and the molecules are incorporated into the growing crystals. In

quiescent solutions, metastable lamellar shaped crystals are formed which are

laterally large but possess a limited thicknesses (50 . 200 A) in the chain direction

due to chain-folding. In stirred solutions, crystallization result:5 in the formation of

so-called shish-kebabs, fibrillar crystals consisting of a (partly) extended-chain

crystalline core surrounded by folded-chain lamellae.

The melting temperature of polymer crystals in semi-dilute solutions can be

approximated by equation 3.1 [1]:

Blending of Trans-l,4-Polybutadiene and UHMW-Polyethy'.!!:e in Solution 37

(3.1)

where T m 0 "" the equilibrium melting temperature of the pure polymer, T m ... the

equilibrium dissolution temperature of the polymer in the solvent V u '" the molar

volume of a monomer unit, (/)1 '" the volume fraction of solvent, t..Hu == the melting

enthalpy per mole monomer units, R ;;;;; the universal gas constant, and X = the

interaction parameter for polymer and solvent.

Crystallization of polymers from solution is usually referred to as liquid-solid

deruixing (L-S). Besides liquid-solid demixing, polymer-solvent systems can also

show liquid-liquid (L-L) demixing. The free energy of mixing of a polymer-solvent

mixture can be described by equation 3.2, which Was developed independently by

Flory [2,3], Huggins [4,5] aIJd StaverroaIJ [6].

(3.2)

where AG is the free energy of mixing per mole lattice sites, m1 is the relative chain

length of the solvent molecules, ~ is the relative chain length of the polymer

molecules and t/)2 is the volume fraction polymer.

Comparing equations 3.2 and equation 3.1 shows that liquid-liquid and liquid-solid

demixing are influenced by the same parameters and it is possible that both

phenomena occur in the same temperature region.

The possible interferences ofliquid-liquid and liquid-solid (crystalli:lation) dt:mixing

for polymer-solvent mixtures are illustrated schematically in figure 3.1. In this fi~'U .. e

the liquid-liquid demixing curve for systems with an upper critical solution

38 Chapter 3

temperature (UeST) is combined with a melting/dissolution curve and a crystal.

lization curve. The crystallization curves are dependent on kinetic paramete(s like

the cooling rate.

1. Figure 3.1 a shows the melting- and crystallization curves in the case that no

demix.ing occurs. This kind of behaviour is observed for the system

polyethylene/xylene [7,8]. Upon cooling the homogeneous solution, crystal­

lization takes place and a gel-like polymer network is formed. In very dilute

solutions no network is formed but separate single crystals due to a lack of

overlap between the different molecules.

2. The L-L demixing curve is located below both the crystallization and the

dissolution curve (figure 3.1b). For some polymer-solvent mixtures, e_g.

isotactic polystyrene in cis·decaline [9]. a flattening of the T m - 4'2 and the

Tc - 1./12 curves can be observed, even if the demixing curve is situated far

below these curves.

3. The L-L dembdng curve intersects the crystallization curve over a limited

concentration range but does not intersect the melting/dissolution curve (see

figure 3.1c). Upon cooling a solution with a concentration within the L-L

demixing range, de mixing occurs, followed by crystallization. The resulting

morphology can be quite complex depending on whether spinodal or binoda!

demixing occurs, cooling rate etc .. The L-L demixing becomes noticeable in

a morc or less pronounced flattening of the melting curve. This kind of

behaviour was observed for the system polyvinylalcohol/ethyleneglycol

[10].

4. The L-L Jemixing curve intersects both the crystallization and the dissolution

curves (see figure 3.1d). This results in a more or less constant crystallization

and melting temperature in the L·L de mixing range as was found, e.g., for

polyethylene in oi-phenyl-ether or in amyl acetate [7,8J

Blending of Trans-J,4-Polyburadiene and UHMW.Polyethylene in Solution 39

T T

.. --J"--....... --,... ... - -.-----.-

,.---.--.. ,.,,"

o o

T T

-~~-------~-~~ ....... -'.~ .................................. .

o o

Figure 3.1 Schematic representation of the influence of liquid-liquid demixing on

the melting and crysrallization curveS of polymeNolvenr systems

(-- L-L demixing, ------ melting, -,-",. crystallizarion)

40 Chapter 3

3.3 Experimental

3.1 Solution.Crystallization of t·l,4-PB, UHMW-PE and their Blends from

Xylene

The materials used in this study werC UHMW-PE, Hostalen Gur-412 supplied by

Hoechst Ruhrchemie (Mw = 1500 kg/mole, Mil = 200 kg/mole) and trans-l,4"poly­

butadiene, synthesized in Our laboratory (M.. = 75 kg/mole, vinyl content

0_9.1.2 %).

Solutions of pure t-1,4-PB in xylene were prepared with several concentrations

(w/w). The crystallization temperatures of t·l,4-PB in these solutions were

determined by differential scanning calorimetry using Perkin Elmt:r large volume

capsules (75 I-d) which, if sealed properly, can sustain pressures up to 24 atm. In the

DSC. The solutions were heated to 140°C and kept at that temperature for 15 min.

Subsequently, the solution was cooled at a rate of 10°C/min to -100°C. In a

subsequent scan the heating CurvC was measured, also at a rate of 10 °C/min, to

determine the melting/dissolution temperatures_

The possible occurrence of liquid-liquid phase separation in the system t-1,4"

PB/xylene was investigated using light scattering. Solutions with several

concentrations were cooled from 120 QC to 10°C in a closed cell with a transparent

top and bottom- A Ia.~er beam (488 nm) was passed through the cell and a detector,

placed under an angle of about 45 ~ to the laser beam, was used to detect scattered

light due to potentialturbidiry occming in the solution as a result of L .. L demixing

or crystallization. The cell was placed in a heating device which permitted

controlled heating and cooling at a rate of 1 "C/mjn. The temperature of the cell

and the intensity of the scattered light were recorded simultaneously_

In the same way, the occurence of L.-L demlxing in the ternary system UHMW­

PE/t-l,4-PB/xylene was investigated using a solution containing 15 g polymer per

1000 ml xylene (the concentration used for film casting, see section 3.3,2)_ Because

blends are studied with several t· l,4-PB contents, the ratio UHMW-PE!t-1,4-PB in

this solution was varied.

Blending of Trans-l,4-Polybutadiene and UHMW-Polyethylene in Solution 41

The crystallization conditions used in the DSC experiments (cooling rates of

10 "C/min) simulate the crystallization conditions occuring during the film casting

procedure described in section 3.3.2. Due to the relatively large sample volume used

for the light scattering experiments, the cooling rate in these measurements was

limited to only 1 "C/min.

3.3.2 Preparation of Blended Films

Blended films of t-I,4-PB and UHMW-PE were prepared by solution crystallization_

Both components were mixed in several ratios and 0.5 % (w/w) stabilizer di-t-butyl­

p-cresol (DBPC) wa, .. added. The materials were suspended in xylene up to a

concentration of 1.5 % (w Iv). The obtained suspension was degassed under vacuum,

then saturated with nitrogen gas and subsequently heated in a silicone oil bath to

approx. 120°C. During heating the suspension was stirred to a homogeneous

dispen;jon. The stirring was stopped as soon as dissolution started. This could be

observed from agglomeration of powder particles and the "Weissenberg effect"

around the stirrer. The dispersion was kept at 120 ac for 2 - 5 hours to obtain a

(macroscopically) homogeneous solution. This solution was poured into an

aluminum tray. Upon cooling, crystallization/gelation occurred and the solvent was

slowly evaporated in a fume cupboard at room temperature. The resulting films

were extracted with n-hexane at room temperature to remove the stabilizer_ After

extraction, the films were pressed at room temperature to obtain fiat and void free

samples.

3.3.3 Differential Scanning Calorimetry (DSC)

Thermograms of the blended films were recorded using a Perkin-Elmer DSC-J

differential scanning calorimeter. A standard heating rate of 10 °C/min was

adopted. The temperatures reported in table 3.1 are the peak temperatures of the

42 Chapter 3

transition under investigation.

Indium was used for temperature calibration (Tm '" 156.6 T, AHrll "" 28.4 Jig).

3.3.4 Transmission Electron Microscopy (TEM)

Transmission Electron Microscopy was performed using a Jeol JEM 2000 FX

microscope, operated at 80 kV. Samples were treated with a 2 % Os04-solution at

room temperature during 24 hours and then embedded in an epoxy matrix.

After embedding, the specimens were trimmed and subsequently treated during 16

hours with a Ru04-solution prepared according to Monte~inos et a1.[l1]. Finally

thin sections were obtained by ultramicrotomy at room temperature using a

Reichert Ultracut E microtome.

3.3.5 Wide Angle X-ray DitTraction (WAXD)

Wide Angle X-ray Diffraction (W AXD) patterns were obtained using a Statton

camera_ Ni-filtered Cu K.:t-radiation was generated at 50 kV and 30 rnA.

3.3.6 Fourier Transform Infra-Red Spectroscopy (FI'IR)

Infra-red spectra were recorded with a Mattson Polaris spectrometer equipped with

a standard DTGS detector and a HeiNe laser. The spectra were obtained after

accumulating 32 scans at a resolution of 4 cm-! between 4000 and 400 cm-1• The

blended films were placed directly in the spectrometer, without further preparation_

The laser beam was directed perpendicular to the film surface. Measurement~ were

performed at room temperature_

Blending of Trans-l,4-Polybutadiene and UHMW-Polyelhylene in Solution 43

3.3.7 Quantitative Analysis of the Crystallinity

A quantitative analysis of the crystalline phases of polybutadiene and polyethylene

in the blends is not possible by DSC alone since the melting peaks of t-I,4-PB and

UHMW-PE (partly) coincide. Therefore, a FfIR subtraction method, introduced

by Wang and Woodward [12], was used to determine the crystallinity of t-l,4-PB.

In this method the FTIR spectrum of a semi-crystalline sample is recorded at a

temperature below the melting point and at a temperature above the melting point.

Subsequently, the latter spectrum is subtracted from the first in such a way that a

characteristic amorphous band of the sample completely disappears in the

difference spectrum. The subtraction factor necessary to achieve this goal, is a djrect

measure for the fraction amorphous in the sample.

In the case of t~1,4"PB the non"crystalline bands at 1352 cm· l and 1076 cm- l can be

used for this analysis [13]. In the blends, the 1352 cm-1 is partly overlapped by an

intensive polyethylene band and therefore the 1076 cm"\ has to be used. The ana­

lysis described above has been performed on a blend containing 10 wt.% t-l,4.PB.

3.3.8 Nuclear Magnetic Resonance (NMR)

A blended film containing 10 wt.% t-l,4-PB was investigated by 13C Nuclear

Magnetic Resonance (NMR) at the Institut fUr Makromolekulare Chemie of the

Albert-Ludwigs-Universlit in Freiburg. Measurements were performed below and

above the solid-solid phase transition temperature and a comparison was made with

100 % t-l,4-PB samples. Both 13C-CP-MAS NMR spectra and 13C-DD-MAS side

band spectra were recorded. For details concerning the experimental set up and

conditions we refer to reference [13]. The chemical shifts reported are in ppm's

downfie1d from tctramethylsilane (TMS).

44 Chapter 3 -_ ...... _._---------

3.4 Results and Dis(llssion

3.4.1 Crystallization of UHMW-PE from Xylene

Figure 32 shows the onset temperatures of the crystallization peaks and the final

temperatures of the melting peaks of UHMW"PE (Hostalen Gur 412) in xylene as

a function of concentration. The data [14] are obtained via DSC cooling and

heating scans at a scanning rate of 10 °C/min.

At a concentration of about 1 to 2 wt% (the concentration used for preparing the

blended films) the cl)'stallization temperature of UHMW·PB is about 80 ~c.

2 III 5 .-('(l '--Q)

g-(j) ...,

Fi!:.>'ure 3.2

150 150

125 ."". ....... "

125

• .... " ... /"

ft//"

100 "," 100 .... ~ ~ ____ e

p" tfr

75 75

50 50 0-00 0_20 0.40 OBO 0.80 1.00

fraction PE (-)

Final melting temperatures (_) and onset temperatures of crystallization

(0) of UHMW·PE in xylene as a /unction of concentration

Blending of Trans-l,4-Polybutadiene and UHMW-Pofyetlzylene in Solution 45

3.4.2 Crystallization of T -I,4-PB from Xylene

DSC cooling curves of the t-l,4-PB solutions in xylene show one peak at relatively

low temperatures due to crystallization. During subsequent heating, two partly

overlapping peaks are found. We assume that the first peak is due to the transition

of the monoclinic crystal structure of t-l,4-PB into the hexagonal crystal structure

and that the second peak is due to dissolution of the latter structure.

The final temperature of both transitions of the double peak observed in the DSC

heating cuJVes are shown in figure 3.3 as a function of the polymer concentration

in solution. It appears that the curve related to the first transition peaks observed

upon heating (the monoclinic into hexagonal transition), shows the curvature

expected for a 'normal' melting point depression curve, the shape is comparable to

the melting curve of polyethylene in xylene (see figure 3.2).

The melting temperature of the hexagonal t-l,4-PB crystals On the contrary, shows

an extremely large decrease (almost 90 0c), even when only a smaIl amount of

solvent is present. This decrease is far too large to be explained by a 'normal'

melting point depression. Furthermore, the melting temperature appears to be

constant over a very large concentration range (between 20 and 70 wt.% t.l,4.PB).

An invariable melting temperature over a certain concentration range is indicative

for the occurrence of L-L denrlxing (see figure 3.1d). The flattening of the melting

curve (curve b) observed in the case of t-l,4.PB in xylene however, is spread over

a concentration range which is much broader than usually observed for L-L

demixing in polymer I solvent mixtures. Furthermore, by light scattering experiments

Over the concentration range from 20 to 70 wt.% t-1,4-PB, no cloud points due to

L-L demixing, were observed upon cooling down from 120 "c. Therefore, it is

unlikely that L-L demixing is the cause for the unusual shape of the melting curve.

A possible explanation for the observed phenomena is that the crystals, as soon as

they have transferred into the conformationally disordered hexagonaJ stmcture,

46 Chapter 3

'ab~()rb' ~olvent molecules, thus forming a separate phase which immediately starts

to melt/dissolve.

This explanation suggest that the hexagonal crystal structure is not stable in xylene

solutions, which i$ confirmed by the very low crystallization temperatures of t·'l ,4.PB

in xylene (see culVe 3.3c). It is not dear whether crystallization into the hexagonal

or into the monoclinic crystal structure occurs at the crystallization peak measured

by DSC. However, the temperatures at which crystallization starts, is within the

temperature range where in the Ca$e of pure t-l,4-PB the monOclinic crystal form

is stable, indicating that hexagonal crystals are not formed in the presence of xylene,

Further research will be necessary to understand the peculiar phase behaviour of

the sy~tem t-l,4-PB/xylene in mOre detaiL

2 ~

:J ... ItS iii g 2

Figu.re 3.3

150

130

110

gO

70

50

30

10

-10

_ . .0------~o

-Q~~

I

I )

(

I .~

150

130

110

90

70

50

30

10

-10

-30 ~--~--~~--~--~~--~-30 0.00 0.20 OAO 0.60 0.80 1.00

fraction t-1.4-PB (-)

Onset temperatures of crystallization (0), final melting temperatures (D)

and final temperatures of the solid-solid-transition (_) of t-l,4-PB in

xylene as a .function of concentration

Blending of Trans-l,4-Polybutadiene and UHMW-Polyethylene in Sol~~~,! 47

The reduction in crystallization temperature of t-I,4-PB, due to the addition of

xylene, appears to be much larger than in case of UHMW.PE. In a solution

containing 1 wL% t-l,4-PB, the crystallization temperature is about 4 OC whereas,

in a comparable solution, UHMW-PE shows a crystallization temperature of about

80°C.

In practice, for the preparation of the blends from xylene solutions, this difference

in crystallization behaviour of UHMW-PE and t-l,4-PB from dilute xylene solutions

implies, that solution blending of these two polymers and subsequent cooling will

result in phase separation (L-S) induced by crystallization: Upon cooling a solution

containg both PE and t-l,4-PB, first UHMW-PE is expected to crystallize at about

80 <oC, forming a gel swollen with xylene in which t-l,4-PB is still dissolved. Either

further cooling to below 4 ~C would be necessary to crystallize this t-l,4-PB or

evaporation of the solvent, which increases the t-l,4-PB concentration and thus the

crystallization temperature.

DSC experiments on mixed solutions containing 15 g polymer (UHMW.PE and

t-l,4-PB in several ratios) per 1000 ml xylene indeed reveal that subsequent

crystallization of both components occurs upon cooling the mixed solutions (see

table 3.1).

Interesting is that, although the crystallization temperature of UHMW.PE in the

mixed solutions hardly changes compared to the pure UHMW-PE solution, the

crysta11i~ation temperature of t·l,4-PB increases with an increasing fraction of

UHMW.PE present in the rllixed solutions. This can be explained by assuming that

the polyethylene crystals in the swollen gel act as nuclei for the crystallization of the

t-l,4~PB: the higher the ratio PE/PB, the more polyethylene crystal surface is

available to act as a nucleus for the crystallization of t-l,4-PB.

4X Chapter 3 --_________ u •• _.,,···-_______ _

Table 3.1 Crystallization temperatures of UHMW"PE and t-l,4-PB in xylene

solutions containing 1.5 % (wjv) polymer (UHMW-PE and t-I,4-PB in

s(~veral ratios)

ratio PE/PB Tc.unset(PE) lc,cm .. t(PB)

(w/w) CC) eq

100/0 81

95/5 83 27

80/20 84 26

65/35 84 15

50/50 84 17

0/100 5

3.4.3 The MOl'Phology of Blended Films

Preparation

The preparation of the blended films was described in section 3.3.2. After mixing

UHMW-PE and t·J,4.PB in solution at 120°C, the solution is poured into a tray

and left to cool to room temperature_ Subsequently, the xylene is allowed to

evaporate_ Concerning the cooling conditions, this procedure resembles the DSC

experiments described in the previous section. However, the final temperature

(room temperature) is for some of the PE/PB mixtures in xylene above the

crystallization temperature of t-l,4-PB as determined by DSC (see table 3_1).

Therefore, it i~ likely that in these cases the t-l,4-PB present in the gel will only

start to crystallize when, through evaporation of xylene, the concentration has

increa<;ed.

Blending of Trans-I,4-Polybutadiene and UHMW-Polyethylene in Solution 49

Microscopy

The morphology of the dried blended films is studied using TEM. Transmission

electron micrographs of the solution-crystallized blends clearly show the presence

of both t-l,4-PB and UHMW-PE lameIJae, the former ones being darker (due to

specific staining with OS04) and thicker than the latter (approx. 125 and 70 A resp.,

see figure 3.4.a). Both kinds of lamellae show the same preferential orientation

parallel to the film surface.

With increasing t-l,4-PB content, besides the individual t-l,4-PB lamellae (clearly

visible in figure 3.4a) sometimes larger and more irregularly shaped dark domains

are visible in the micrographs (see e.g. fig 3.4c). This seem to be agglomerates of

t-1,4-PB in which no stacks of lameIJae can be discerned which do contain

crystalline regions (see figure 3.5).

Figure 3.5 TEM micrograph of a t-I,4-PB domain in a blend containing 20 wt. %

t-I,4-PB

The micrographs clearly show that areas of both mainly UHMW-PE and mainly

t-l,4-PB exist in the solution-crystallized blends. This is in agreement with the

crystallization behaviour of both components in xylene as discussed in the previous

section. UHMW-PE and t-l,4-PB crystallize one after the other, UHMW-PE upon

cooling to about 80 ¢C and t-l,4-PB later, after evaporation of a part of the xylene.

50

Figure 3.4

Chapter 3

Transmission electron micrographs of a solution-crystallized blend film

containing 10 wt.% t-1,4-PB (a), 20 wt.% t-1,4-PB (b) and 50 wt.%

t-1,4-PB (c)

Blending of Trans-l,4-Polybutadiene and UHMW-Polyethylene in Solution 51

The existence of separate domains of each constituent is confirmed by X-ray

diffraction and Ff-IR experiments which both show the characteristics of the pure

components. Figure 3.6 shows the W AXD pattern of a solution-crystallized blend

containing 33 wt.% t-1,4-PB. Two intense reflections can be observed, assigned to

the orthorhombic (110) and (200) lattice planes of polyethylene. Between these two

intense PE reflections, a relatively weak t-l,4-PB reflection is visible. This reflection

(2 e = 22,4°) has been assigned to the (120) and (200) monoclinic lattice planes

[15].

Figure 3.6 W AXD pattern of a solution-crystallized blended film containing

33 wt. % t-l,4-PB (X-ray beam paralleL to film surface)

It should be noted that the reflection rings of both components show intensity

maxima at the poles caused by the preferential crystal orientation parallel to the

film surface, as discussed before.

The infra-red spectra of the blends contain the characteristic bands of both

UHMW-PE and t-l,4-PB. Figure 3.7 shows the IR spectrum of a solution­

crystallized blend containing 20 wt.% t-l,4-PB and of a comparable 100 % UHMW­

PE film. The additional band at 970 cm· l caused by the C = C bonds in t-l,4-PB

52 Chapter 3 --------_ ..... _ .... __ ._-_ .. _-

is clearly visihle just like the characteristic t-1 ,4-PB bands corresponding to the low

temperature monoclinic lattice, i_co, 1236 cm-], 1054 cm- l and 772 cm- l [16,17}.

Fif.,'Ure 3.7

1400 1200 1000 IlOO

wavenumbers (om-I)

FT-JR spectra of a solurion-crystallized blended film containing 20 wt. %

t.J,4.PB and of comparable 100 % UHMW-PE film

Blending the two components does not induce shifts in these characteristic

crystalline t-l,4-PB bands, again indicating that separate crystallization of both

constituents has occurred.

Quantitative analysis of the crystallinity

In figure 3.8, DSC thermograms are given of three solution-crystallized UHMW­

PR/t-1,4-PB blends with different compositions_ The thermograms show the melting

endotherms of both UHMW-PE and t-l,4-PB. In the blends with relatively low t·

l,4-PB contents, the melting peak of t-l,4.PB is difficult to discern because it

(partly) coincides with the UHMW-PE melting peak. However, a small shoulder is

visible on the left hand side of the UHMW-PE melting peak at about 128 °c, probably due to the melting of t-1,4-PB..

Blending of Tralls-I,4-Polybutadiene and UHMW.Polyethylene in Solution 53

50 1 0 tempen!,lure (DC)

Figure 3.8 DSC thermograms of solution"crystallized blended films containing (0)

10 wt.%, (6) 25 wt.% and (e) 50 wt.% t-l~4-PB

The thermograms also show the solid·solid phase transition endothenn of t-1,4-PB

around 60 - 70 "C, which consists of more than one peak. This phenomenon has

been observed before upon crystallization of t-l,4-PB from solution and it is

contributed to the presence of monoclinic crystals of different thermal stability

[18,19,20,21].

In table 3.2 the measured (combined UHMW-PE and t-1,4-PB) melting enthalpy

and transition enthalpy of several blends are given. The ratio of the measured

melting enthalpy (.6.HI1l ,01<p) over the theoretically expected melting enthalpy (.6.Hm ,th) •

.6.Hm•1h is a qualitative measure for the overall crystallinity of the blends . .6.Hm.1h was

calculated from the blends composition and the literature values for the melting

enthalpy's of 100 % crystalline UHMW-PE and t-l,4-PB (293 and 69 J!g respec­

tively) using equation 3.3.

54 Chapter 3 ----......... _--

in which WJ'fl and wPB represent the weight fractions of polyethylene and

polybutadiene in the blend~.

Table 3.2 Quantitative anaiysl.r of the crystalline phases in UHMW-PE/t-I;4-PB

blends

PE/PB ABm,lb 4Hm , • .-p 11 Hexp mo~_hex

(-) (J/g) (J/g) (J/g)

90/10 271 198 50

80/20 248 165 53

75/25 237 163 59

50/50 181 148 73

Comparing the solid-solid phase transition enthalpy measured by DSC

(I1Wxp mon_hex) with the expected value for 100 % crystalline t-l,4-PB (144 J/g)

reveals a large discrepancy between these two values (see table 3.2). Since we know

that co-cry~tallization of PE and t-l,4-PB does not occur, thi~ discrepancy can be

attributed to a low crystallinity of the t-l,4-PB domains or to a fraction t-1,4-PB

which is already at room temperature in the hexagonal phase, or to a combination

of these two factors.

To gain some more in~ight in the crystalline phases of the t-l,4.PB domains in the

blends, the crystallinity of t-l,4,PB in one blend (containing 10 wt.% t"1,4-PB) was

studied in more detail, via the FT-IR subtraction method described in section 3.3.7.

~lending ;:f TranS-l,4-Polyhutadiene and UHMW-Polyethylene in Solution 55

The IR spectrum of the molten blend (recorded at 150 cC) was subtracted from the

spectrum of the blend at room temperature, in such a way that the intensity of the

1076 cm-1 band (representative for the amorphous fraction in t-1,4-PB) became O.

The factor necessary for this substraction operation was 0.5 ± 0.1 indicating that the

crystallinity of t-l,4-PB in the 10 % blend amounts 50 % (± 10 %).

Taking into account a crystallinity of 50 %, the fraction monoclinic of the crystalline

t-l,4-PB domains can be calculated to be 70 %, indicating that 30% of the

crystalline t-l,4-PB fraction in the blend is jn the hexagonal phase at room

temperature.

Considering a crystallinity of t-l,4-PB of 50 %, the crystallinity of PE can be

calculated from the combined UHMW-PE/t-l,4-PB melting enthalpy as measured

by DSe and it appears to be about 75 % which is consistent with values reported

for pure UHMW-PE films prepared under similar conditions.

Summari:dng: a blend containing 90 wt.% UHMW-PE and 10 wt.% t-l,4-PB

contains PE with a crystallinity of 75% and t-1,4-PB divided in three fractions, 50%

amorphous, 35 % monoclinic and 15 % hexagonaL

The presence of a hexagonal t-1,4-PB fraction in the blends at room temperature

could not be observed by WAXD but this is due to the fact that the corresponding

reflection (29 "" 20.8 ~) is very weak compared to the other reflections and partly

coincides with one of the polyethylene reflections. He NMR measurements however

do reveal a hexagonal t-l,4-PB fraction existing at room temperature.

The Be NMR spectrum of a blend containing 10 % t-l,4-PB is given in figure 3.9.

The resonances around 131 ppm are due to the trans olefinic C atoms in t-l,4~PB.

56

140

Figure 3.9

__ , .. _Chapter 3

35_3 131.8

ssb

100 60 20 chemical shift (ppm)

l.lC NMR spectrum of a solution-crystallized blended film containing 10

wt.% t-l,4-PB

TIle trans olefinic signal consists of several resonanceS. Usually the resOnance at

130_6 ppm is assigned to the amorphous and/or hexagonal phase and the resonance

at 131.8 ppm to the monoclinic phase. It is not possible to discriminate between an

amorphous and a hexagonal phase in these spectra. In principle, an analysis of the

methylene resonances could solve this problem [22] but in the blends an intensive

and broad PE resonance at 35.3 ppm, overlapping the methylene resona.nce of t-1,4-

PB, makes this procedure impossible. Therefore, Deckmann [13} performed slow­

spinning NMR experiment" to determine the chemical shift anisotropy (CSA) which

can be used to discriminate be teen crystalline a.nd non-crystalline components, It

appears that the resonance at 130.6 ppm is at least partly due to t- t,4-PB in the

hexagonal modification_

Miscibility in the amorphous phase

The analysis concerning the morphology/miscibility of PE/PB blends as presented

so far in this chapter, has been focllsed on the crystalline domains in the blends. It

was shown that co-crystallization of PE and t-I,4-PB does not occur. However,

considering the preparation method of the blends, some potential miscibility of PE

Blending of Trans-I/4-Polybutadiene wzd UHMW-Polyethylene in Solution 57

and PB in the amorphous domains can not be excluded. When UHMW-PE has

crystallized after casting of the SOlution, the amorphous UHMW.PE parts (e.g.

interlamellar tie-molecules, folds and chain-ends) as well as t-l,4-PB remain in

solution. Solidification of these components starts upon evaporation of xylene and

can possibly occur simultaneously, resulting in pardy mixed amorphous domains.

The presence of a small amount of t-l,4-PB chains in the amorphous PE domains

'Will be difficult to detect. DMTA experiments were performed to obtain the tano

curves of a 100 % UHMW-PE film and of a blend film containing 20wt.% t-l,4-PB.

The results of these measurements however, are not conclusive. Similarly, the FT­

IR data give no proof for partial miscibility of both components in the amorphous

phase, however, due to the absence of strong specific interactions between PE and

t-l,4-PB no significant shifts in the amorphous bands in the IR spectra can be

expected. Certainly not, if only a small amount of t-l,4-PB is present in the

amorphous PE domains.

In conclusion, some t-l,4-PB chains might be trappt:d within the amorphous

UHMW-PE domains, but this can not be demonstrated by the techniques used. The

possiblity of such partial miscibility in the amorphous domains will be addressed in

chapter 6.

3.5 Conclusions

Crystallization of UHMW-PE and t.l,4-PB in xylene at an overall concentration of

appr. I to 2 wt.%, takes place at crystallization temperatures of approximately 80 °c and 4 .. c respectively. A possible explanation for the relatively low crystallization

temperature of t-l,4-PB is that hexagonal t-l,4-PB crystals are not stable in xylene.

Crystallization of a mixture of UHMW-PE and t-l,4-PB in xylene, takes place

consecutively, L-S phase separation through crystallization. The crystallization

5H Chapter 3

temperature of UI-IMW-PE is not influenced, but the crystallization temperature

of t-l,4-PB increases with increasing UHMW.PE content in the $ystcm. Probably,

the UHMW-PE lamellae act as nuclei for the crystallization of t-l,4-PB.

Blended films prepared via solution-crystallization consist of separate UHMW·PE

and t-l,4-PB crystals: co-crystallization does not occur. The crystallinity of the PE

domains in the blended films is about 80 %, comparable to the crystallinity of

solution-crystallized pure UHMW-PE films, but the crystallinity of t-l,4-PB is

reatively low, about 50 %. The crystalline t-l,4-PB is partly prescnt in the hexagonal

phase.

In blends with low t·1,4-PB contents « 10 wt.%) t-l,4-PB is present as individual

lamellae in between stacked PE lamellae, at higher t-l,4-PB contents agglomerates

of t-1,4-PB crystals are formed_

3.6 References

1. Wunderlich, 8., 'Macromolecular Physics, Volume 3: Crystal Melting',

Academic Press, New York, 1980, pA2

2_ Flory, PJ., J- Chern. Phys. 1941,9, 660

3. Flory, P.J., 1. Chern. Phys- 1942, 10, 51

4_ Huggins, M.L, Ann. N.- Y. Acad. Sci. 1942, 43, 1

5. Huggins, M_L,1. Chern. Phys. 1941, 9, 440

6- Staverman, AJ., Rec- Tr(I1J_ Chim. 1941, 60, 640

7. Richards, RB., Trans. Faraday Soc. 1946,42, 10

8. Flory, P.1_ and Mandelkern, L,l. Am. Chern. Soc- 1951, 73, 2532

9. Jacobs, A and Berghmans, H., Eur_ Phys. Con! 1989, 13F, 50

10_ Berghmans, H. and Stob, W., 'Integration of Fundamental Polymer Science

and Technology', Ed. L.A Kleintje5 and PJ. Lemstra, Elsevier AppL Sci.

Pub!', London, 1986, p_ 218

Blending oj Trans.l,4.Polybutadiene and UHMW-Polyethylene in Solution 59

1 L Montezinos, D., Wells, E.G. and Burns, J.L.,l. Polym. Sci. Polym. Lett.

Ed. 1985,23,421

12. Wang, P. and Woodward, A.E., MakromoL Chern. 1989, 190,875

13. Deckmann, H., PhD. The.~is, AJbert-Ludwigs-Universitiit, Freiburg,

Germany, 1991, chapter 4

14. M61ler, M., private communication

15. Iwayanagi, S., Sakurai, I., Sakurai, T. and Seta, T., Repts. Prog,-. Polym.

Phys. Japan 1967, 10, 167; J. MakromoL Sci. Phys. 1968, B2, 163

16. Nikitin, V.N., Volkova, LA., Mikhailova, N.V. and Baklagina, Iu.G.,

VysokomoL Soedi~ 1959, 1, 1094

17. Hendrix, C, Whiting, DA and Woodward, AE., Macromolecules 1971,4,

571

18. Tatsumi, T., Fukushima., T., Imada, K and Takayanagi, M., J. Macromol.

Sci. Phys. 1967, Rl, 459

19. Takayanagi, M., Imada, K.., Nagai, A, Tatsumi, T. and Matsuo, T., J.

Polym. Sci. Part C 1967, 16, 867

20. Marchetti, A. and Martuscelli, E., 1. Polym. Sci. Polym. Phys. Ed. 1976, 14,

323

21. Tseng, S. and Woodward, A.E., Macromolecules 1982, IS, 343

22. Moller, M., Makrornol. Chern. Rapid Comrn. 1988,9, 107

Morphology and P!!!perties of Drawn UHMW"PEjT-l,4-PB Ble~ds

Chapter 4

Morphology and Properties of Drawn

UHMW-PE/T .. 1,4 .. PB Blends

4.1 Introduction

61

In chapter 3 it was shown that films prepared from UHMW-PE and t-1,4-PB consist

of a polyethylene matrix with dispersed t-1,4-PB lamellar crystals. The polyethylene

matrix is highly crystalline and consists of stacked lamellae oriented with the chain

axes perpendicular to the film surface. The t-l,4-PB domains, individual lamellar

crystals at low t-l,4-PB contents and agglomerates of crystals at higher t-l,4-PB

contents, are located in between the PE lamellae and are only appro 50%

crystalline.

The blended films were prepared as a precursor for tensile drawing, to obtain

oriented tapes. The deformation mechanism during drawing of solution-crystallized

100 % UHMW-PE films has been studied by van Aerle [1]- He proposed a three

stage deformation mechanism, resulting in a fibrillar oriented structure consisting

of more or less, aligned and extended PE chains.

Tensile drawing of pure t-l,4-PB is not feasible, unless the films are slightly

crosslinked (see section 2.4.3).ln that case, the t-l,4-PB chains are oriented towards

the drawing direction and partly extended (the macroscopic draw ratio is about 20).

62 Chapter 4 ----_._ ........... _ ... _---------- --------

In this chapter the tensile drawing of Solution-crystallized blends of UHMW-PE and

kl,4-PB will be discussed.

4.2 Experimental

4.2.1 Sample Preparation

The materials used in this study were UHMW-PE, Hostalen Gur-412 supplied by

Hoechst Ruhrchemie (M", "" 1500 kg/mole, Mft '" 200 kg/mole) and trans-l,4-

polybutadiene, synthesized in our laboratory (M.. ;;;; 75 kg/mole, vinyl content 0.9-

1.2 %). The hlends were prepared via mixing in xylene and subsequent solution­

crystallization according to the method described in detail in chapter 3. The

resulting films were cut into tapes of 25 x 8 mm2 and drawn manually to several

draw ratios on a hot-shoe, using ink·marks to determine the displacement. The

drawing temperature was approximately 120 QC which is above the solid-solid phase

transition temperature and below the melting temperature of t··l,4.PB as well as

below the melting temperature of polyethylene. At this temperature, pure solution­

crystallized UHMW·PE films can be drawn easily to high draw ratios.

J n general, blends with t -1,4-PB contents up to 20 wt. % were studied except in some

morphology studies (TEM and WAXD) which required a higher t-1,4.PB content

to be able to detect the t-1,4-PB phase.

4.2.2 Morphology

The morphology of the drawn, blended tapes was characterized using DSC, TEM,

WAXD, FT·IR and DC NMR. The experimental parameters and procedures are

described in detail in section 3.3 of the previous chapter.

Morphology and Propertie.s of Drawn UHMW-P£jT.l,4-PB Blends 63

4.2.3 Degree of Orientation

Absorption spectra were obtained in the 4000 - 400 cm-1 region with a resolution

of 4 cm-1 using a Mattson Polaris FT-IR spectrometer equipped with a standard

DTGS detector and HeiNe laser. The absOJptions in the direction parallel (AI)

and in the direction perpendicular to the drawing direction (A..) were measured

using a Specac polarizer consisting of 0.2 /Lm wide aluminum strips on a KRS-5

substrate. From these absorptions the dichroic ratio (D) was calculated using

equation 4.1:

(4.1)

From. the dichroic ratios for the crystalline polyethylene bands at 730 cm-1 and 720

cm-\ the orientation functions for the a-axis (f.) and b-axis (ft» of the polyethylene

crystal were derived, using equations 4.2 and 43:

(4.2)

(4.3)

Since for orthorhombic polyethylene the a, band c axes are mutually perpendicular,

the orientation functions of the three crystal axes are related by equation 4.4;

64 Chapter 4

(4.4)

Consequently, by determining fa and fb via dichroic ratio measurements, the orienta­

tion function of the c-axis of the polyethylene crystal can be calculated.

4.2.4 Tensile Testing

Tensile testing was performed at room temperature on a Frank 81565 tensile tester

equipped with an extensometer. To avoid slippage in the damps, the tapes were,

at both ends, glued in between cardboard tabs.

The cross-sectional area of the tapes was determined from their length and weight

assuming a density of 0.98 gjcm3. The strain rate was l~ S·I with an initial length

between clamps of 150 rom.

The reported values for Young's modulus and tensile strength were averaged over

5 experiments.

4.3 Results and Discussion

4.3.1 nrawability

The mechanical properties of UHMW-PE/t-l,4-PB blends were studied for blends

containing up to 20 wt. % t" 1,4·PB. These blends can be deformed in the solid state,

via uniaxial tensile drawing, up to macroscopic draw ratios of about 170 (see figure

4.1). TIle morphology of the blends after drawing and the resulting properties are

discussed in the following sections.

MOrphology and Properties 0t Drawn UHl:JW-PE/T·l,4-PB B1e..n_cL_~ _____ 6_5

I o ~ (Il .....

Figure 4.1

180 r-----~---------_____,

j

80 ~--~--~--~---~--~ o 5 10 15 20 25

t-1.4-PB content (wt.%)

Maximum attainable draw ratio via single stage, uniaxial tensile

drowing aJ 120 QC, of UHMW-PE/t-l,4-PB blends, as a function of the

t-l,4-PB content

4.3.2 Morphology

Visualization

Upon drawing, drastic changes occur in the morphology of the blends. This is

clearly illustrated by figure 4.2 which shows transmission electron micrographs of

a blend containing 50% t-l ,4-PB after drawing. The micrograph of the undra\m

blend is discussed in chapter 3 (figure 3.4). The micrograph of the drawn blend

shows small stained areas which probably are the t-l,4-PB domains since these react

mOre easily with the staining agents than the PE matrix, due to the unsaturated C-C

bonds and the lower crystallinity. The t-l,4-PB domains are oriented parallel to the

drawing direction. A comparison between figures 4.2a and 4.2b shows that the t-l,4-

PB domains are elongated upon drawing. To what dimensions this thinning process

66 Chapter 4

proceeds upon drawing to higher -draw ratios is not clear, since samples with higher

draw ratios were difficult to stain and section due to the high degree of crystallinity

and orientation.

Figure 4.2 TEM micrographs of drawn UHMW-PE/t-I,4-PB blends containing 50

wt. % t-I,4-PB; a. A,-::::IO, b. 1::::25 (drawing direction is vertical)

The diameter of the stained areas in the drawn 50 wt. % blends at }. = 25, is in the

order of a few nanometres, only 10 to 20 t-l,4-PB chains. This might even be less

in the oriented samples containing < 50 wt.% t-l ,4-PB, since in these blends the t-

1,4-PB domains in the undrawn films are thinner to start with (individual lamellae)

than the agglomerates of t-l,4-PB crystals in the starting material of the 50 wt. %

blends.

Orientation of the chains

The molecular orientation of the UHMW-PE and t-l,4-PB chains after drawing was

studied by wide-angle X-ray diffraction. Figure 4.3 shows WAXD patterns of a

blend containing 33 % t-l,4-PB at several draw ratios. The W AXD pattern of the

undrawn blend (figure 4.3a) shows two intense reflections assigned to the

orthorhombic (110) and (200) lattice planes of polyethylene. In between these two

intense PE reflections a relatively weak t-l,4-PB reflection is visible. This reflection

Morphology and Properties of Drawn UHMW-PE/T-l,4-PB Blends 67

(29 = 22.4°) has been assigned to the combined (120) and (200) monoclinic lattice

planes of t-l,4-PB [2]. Notice that the UHMW-PE and the t-l,4-PB reflections

show the same increase of intensity at the poles. Due to the solution crystallization

method, the crystals are oriented with the chain axes perpendicular to the film

surface in the as-cast films.

Figure 4.3 WAXD patterns of blends containing 33 wt. % t-l,4-PB at several draw

ratios; }'=1 (a), }'=1l (b) and }.=40 (c) (X-ray beam parallel to film

surface; drawing direction is vertical)

68 _________________________ _ --------________ L_f_!a~~~r4

During the drawing process of the blends, UHMW-PE shows its 'normal' behaviour,

the UHMW·PE chains arC rotated towards the drawing direction as appears from

the contraction of the reflection rings at the equator of the WAXD pattern in figure

4.3b. Similar observations were made by van Aerie and Braam for drawing of pure

UHMW-PE [3,4].

Figure 4,4 revcal& the orientation function of the polyethylene crystal c axis (fe) as

a function of draw ratio for a 100 % UHMW-PE film and for a blended film

containing 20 wt.% t·l,4-PB. The orientation of PE chains appears to be not signifi­

(.:;l.ntly affected by the presence of t"l,4-PB.

I

1.00 r-----------------------------,

0.50 I

I 6 I ,6 , : ... ,

... i..-~---I

........ ---- .... --_ ........ --

0.00 *~-----------------------------------~

, I

,

~ -0.50 ~ ____ -L _____ ~----~ ______ ~ ____ ~

Figure 4.4

o 20 40 60 80 100

draw ratio (-)

Orientation function (jcJ as a function of draw ratio for a blend

containing 20 wt.% t-l,4-PB (A) and a 100 % PE sample (11)

The t-l,4-PB in the blends shows the same orientation behaviour as UHMW-PE

upOn tensile drawing (see figure 4.3b). This is in contra&t to the tensile drawing of

pure t-l,4-PB, which results in premature sample failure at draw ratios around 5

Morphology (1nd ~::)perlies of Drawn UHMW·PEjT·I,4.PB Blends 69

[5]. Apparently, the polyethylene matrix. imposes its deformation on the dispersed

t-l,4-PB domains during drawing.

The t-l,4-PB reflection is relatively weak compared to the PE reflection (see figure

4.3a) and upon drawing it broadens in such a way that the intensity maximum is

hardly detectable in the WAXD patterns at low draw ratios « 15) and seems to

disappear completely at higher draw ratios (see figure 4.3c). The broadening of the

t-l,4-PB reflection is probably caused by the thinning of the t-l,4-PB domains as

observed in transmission electron micrographs of drawn blends (figure 4.2b), but it

is not clear why this reflection disappears from the WAXD patterns at higher draw

ratios. It is possible that the thinning of the t-l,4-PB domains proceeds up to

molecular dimensions, in which case no t-l,4-PB crystals are present anymore but

it is also possible that the crystals transform to the hexagonal structure at high draw

ratios. The presence of the hexagonal phase cannot be ruled out, since the most

intense hexagonal reflection, the (100) reflection at 29 "'20.8¢, partly coincides with

one of the reflections of PE (2e "'20.6QC).

The crystalline pbases

DSC measurements performed On blends drawn to several draw ratios, yield a

similar picture as obtained by X-ray diffraction. Upon drawing, orientation and

chain extension of the PE chains occur, resulting in an increase in melting

temperature as can be seen in figure 4.5, which shows DSC thennograms of a blend

containing 20 wt.% t-l,4-PB at several draw ratios.

The combined UHMW.PE/t-1,4-PB melting endotherm shifts towards higher tem­

peratures and the peak width decreases. The ratio of the experimental and the

theoretically expected heat of fusion (calculated according to equation 3.3),

gradually increases with draw ratio from 66 % in the undrawn blend to 81 % in the

blend with draw ratio 80. Due to different ordinate scales of curves a to d in figure

4.5, the crystallinity is not proportional to the melting peak areas shown.

The t-l,4-PB shoulder of the melting peak decreases with increasing draw ratio and

finally disappears or cannot be separated from the UHMW-PE melting peak.

70 Chapter 4 --------------_ .. _------_._-------------_.!...--

Figure 4.5

_ ..... _ ....... _ .. __ .... _-------,

a

b

c

d

50 1 0

DSC thennograms of an UHMW.PE/t-l,4-PB blend containing

20 wt_% t·I,4·PB at several draw ratios (a. A.-l; b. A.~10; c. /1.=30;

d- )=65)

Upon drawing, also the solid-solid phase transition peak of t-l,4-PB disappears from

the thermograms. In a second heating scan, after melting and recrystallization, the

solid-solid phase transition reappears, but the observed transition enthalpy is too

low considering the fraction t·l,4·PB present in the blend. In chapter 3, this

appeared to be due to a low crystallinity of t"1,4-PB and to the fact that t-l,4-PB

is partly present in the hexagonal phase at rOom temperature. In the drawn blends,

similarly, the disappearence of the crystal transformation peak, might be attributed

to an overall decrease in crystallinity of the t-l,4-PB domains Or to a change in

chain-packing from the monoclinic into the hexagonal crystal structure upon

drawing. Both explanations are in agreement with the disappearance of the mono­

clinic t.l,4.PB reflections from the W AXD patterns upon drawing, as was shown in

figure 4.3.

Morphology and Properties of Drawn UHMW.P!l[!"l,4-PB Blends 71

In order to gain some more insight in the structure of the oriented and elongated

t-l,4-PB domains after drawing, Ff"IR and 13C NMR measurements were

performed.

After drawing, the IR spectra of the blends still ~how bands due to the low

temperature crystal phase of t-l,4-PB, indicating that a part of the t-l,4-PB is

present in the monoclinic crystal structure. This in contrast to the results obtained

by WAXD and DSC. However, the intensity of the crystalline t-l,4-PB bands

decreases with increasing draw ratio as shown in figure 4.6.

'0 t:. t!:! E w

Figure 4.6

1.00!li;;-.... ,. ~&;-.. -, ---------------~---.,

.... " ....

" q, 0.75

0.50

0.25

.... ~. "\

"\ A "\

\ 4 '\

\ ~ 'f!,

'n. {j" , , ... , .... f:,

-8~EI_ B------El----

0.00 L--___ ~_'--___ ~.,.....,..,.J~~ ___ ........,I

25 50 75 100

temperature ("C)

The intensity 01 the crystalline 1054 cm-J band as a function of

temperature, lor t-l,4-PB in an UHMW-PE/t-l,4-PB blend containing

20 wt. % t·l;4·P8; (A) 1. ""-1; (D) ). ""-40

In figure 4.6 the intensity of the crystalline 1054 cm· l band (assigned to the low

temperature crystal form [6,7]) is given as a function of temperature, for a blend

containing 20% t·l,4·PB at two different draw ratios.

72 Chapter 4 ---------------~ ... -~.~ .. " .... ~--.-------------=.--

Apparently, the fraction t-l,4-PB present in the monOclinic phase decreases with

increasing draw ratio but is still detectable by IR 5pectroscopy and not by X-ray

diffraction and DSC.

The solid-solid phase transition temperature :shift:s towards a lower value, which is

in agreement with the decrease in crystal dimensions as suggested by the

broadening of the monOclinic rdlection observed in X-ray experiments (figure 4.3b)

and the thinning of the t-l,4-PB domains observed in TEM micrographs (figure

4.2b).

13C NMR spectra of blends containing 20 % t-I,4.PB at several draw ratios are

given in figure 4.7. The resonances between 128 and 133 ppm are due to the trans

olefinic C-atoms in t-l,4·PB. The signal of the aliphatic C-atoms of t-l,4-PB

unfortunately coincides with the resonance of the aliphatic polyethylene C·atoms

(35.3 ppm) and can therefore not be used to study the phase behaviour of t-1 ,4-PB.

The trans olefinic 5ignal consists of several resonances. Usually the resonance at

130.6 ppm is assigned to the amorphous and/or hexagonal phase (it is not possible

to discriminate between amorphous and hexagonal in these spectra) and the

resonance at 131.2 ppm to the monoclinic phase [8). The ratio between the peak

at 130.6 ppm and the peak at 131.2 ppm can be used to evaluate the presence of

the monoclinic and the hexagonal/amorphous phase.

Comparing the spectra in figure 4.7 reveals that the ratio between monoclinic and

(hexagonal/amorphous) decreases with increasing draw ratio. At a draw ratio of 80

the monOclinic phase is still present in these blends, albeit only to a small degree.

On the contrary, blend~ containing 10 wt.% t-l,4-PB show no resonance::; due to the

monoclinic phase anymore already at a draw ratio of 40 [9]. Only resonances due

to higher mobile hexagonal and amorphous t·l,4-PB chains are found.

Morphology and Propenies of Drawn UHMW-PE/T-l,4·PB Blends 73

Figure 4.7

134 132 130 128 126 chemical shift (ppm)

13C·NMRspeclra of an UHMW·PE/t·l,4·PB blend containing 20 wi. %

t-l,4-PB at several draw ratios (a. ). ;;1; b. ). ;;;10; c. ). ... 40; d. ). =80)

The NMR spectra show that the monoclinic t-l,4·PB fraction decreases gradually

upon drawing which is consistent with the data obtained via FT-IR, WAXD and

DSC According to NMR data from blends with relatively low t-1 ,4-PB contents (10

74 Chapter 4 ----------------------------------------

wt.%), the monoclinic fraction even completely disappear~ leaving the t-IA-PB

chains partly in thc hcxagonal and parlly in the amorphous phase. It is not clear to

what degree the resulting chain5 are amorphous or hexagonal. In either case,

amorphous or hexagonal, the mohility of the chains will increase favouring

crosslinking during electron beam irradiation, see chapler 5.

The changes ohserved upon drawing, from a relatively close-packed monOclinic

structure into a more loosely packed hexagonal or even random amorphous

structure, is rather surprising. In UHMW-PE, crystallinity increases upon drawing

a5 shown by an increase in melting enthalpy and melting temperature with

increasing draw ratio [l0, 11, 12].

A similar increase was observed upon orientation of t-l,4-PB films (see chapter 2,

section 2.4.4). From Be NMR studies on t-l,4·PB tapes obtained via drawing of

slightly irradiated t-l,4-PB films, it appeared that the monoclinic fraction increases

upon drawing at the expen:;e of hexagonal fraction (see also chapter 2, section

2.4.4 ).

The fact that t-l,4-PB in blends with UHMW·PE shows the opposite behaviour

might be caused by the thinning of the crystals upon drawing. A~ shown above, this

results in long, thin t·1,4.PB domains with relatively many chains at the surface of

crystals, which are in direct contact with the surrounding PE chains and probably

loose the monoclinic packing. Only the crystal core, which becomes thinner upon

drawing, remains in the monoclinic crystal structure. In blends with higher 1-"1,4.PB

contents the t" 1,4-PB domains are larger (see chapter 3, figure 3.4) and the

monoclinic core survives up to higher draw ratios_

Summarizing, after drawing the hlended tapes consist of thin, elongated t-l,4-PB

domains (a few nanometers thick) dispersed homogeneously in a PE matrix_ This

matrix consists of extended·chain PE fibrils as in case of 100 % UHMW-PE tapes,

prepared via solution-crystallization and subsequent drawing_ Both the t-lA-PB and

the UHMW-PE crystals are oriented towards the drawing direction.

Morphology and Properties of Drawn UHMW-PE/T-I/4-PB Blends 75

The elongated t-l,4.PB domains consist of a crystalline monoclinic core surrounded

by amorphous and/or hexagonal t·l,4-PB. The amount of material in the monoclinic

core decreases with increasing draw ratio, probably due to a further thinning of the

t-l,4-PB domains. The degree of chain extension of the t-l,4-PB chains in the

elongated domains is not known.

4.3.3 Tensile Properties

In figure 4.8, the Young's modulus and tensile strength of UHMW-PE/t-1,4-PB

blends are given as a function of draw ratio for blends with several compositions.

These graphs show that modulus and tensile strength decrease with increasing t-1,4-

PB content.

Of course, a decrease in modulus and strength is espected since, the t-l,4-PB used

has a much lower molecular weight (influencing tensile strength) [13] and a lower

crystal modulus (in the order of 90 OPa for the monoclirtic lattice, cf. chapter 2 and

[l4D·

At a given draw ratio, the absolute values found for modulus and tensile strength

of the tapes, including the values for the 100 % PE tapes, are relatively low

compared to literature values reported for UHMW-PE fibres (see e.g., [l5D. This

is probably due to a different starting morphology of the undrawn fibres and films

and to differences in drawing efficiency.

76 Chapter 4 ._-_ ............ _ .. _ ... -._---------'---

150 ,-------------- .------,

1 25 50 75 100 125 150

draw ratio (-)

2.50 r-----------~----~..........,

2.00

1.50

1.00

0.50

25 50 75 100 125 150

draw ratio (-)

Figure 4.8 Young's modulus (a) and tensile strength (h) as a function of draw

ratio for UHMW·PE/t-l,4-PB blends with 0 (+),1 (.6.),5 (0),10 (D)

and 20 (<:7) wt. % t.J/4.PB

M01phology and Properties oj Drawn UHMW-PE/T-l,4-PB Blends 77

4.4 Conclusions

Blends of UHMW-PE and t-l,4-PB can be drawn via uniaxial tensile drawing at

120°C to draw ratios up to 170 depending on the t-l,4·PB content of the blends.

Upon drawing, the PE crystals in the blends deform in the same way as encoun­

tered upon drawing of pure UHMW"PE films. The t-l,4-PB crystals show the same

orientation behaviour as the UHMW-PE crystals in the initial stages of drawing.

Further drawing results in elongated t.1,4wPB domains of which only the core is

present in the monclinic phase. Upon further drawing the fraction monoclinic t-l,4-

PB decreases in favour of the hexagonal and/or amorphous phase.

Young's modulus and tensile strength of the blends decrease with increasing t-1,4"

PB content.

Drawing of UHMW-PE/t-l,4-PB blends results in oriented structures with a

relatively high modulus and tenacity (dependent on the t-l,4-PB content), consisting

of finely dispersed PB domains in a fibrillar PE matrix. The t-1,4-PB domains are

elongated to almost molecular dimensions and randomly distributed throughout the

tapes. In view of the abundance of unsaturated C-C bonds in these structures it

seems interesting to study the crosslinking of these drawn blends by elecron beam

irradiation.

4.5 References

1. van Aerie, N.AJ.M., Ph.D. ThesL~, Eindhoven University of Technology, The

Netherlands, 1991

2. Iwayanagi, S., Sakurai, I., Sakurai, T. and Seto, T.,l. Macromol. Sci. Phys.

1968, B2. 163

78 Chapler 4

3. van Aerie, N.AJ.M. and Braam, AW.M., J. Maler, Sci. 1988,23,4429

4. van Aerie, N,AJ.M. and Braam, AW.M., ColI. Po(ym. Sci. 1989,267, 323

5, van Aerie, N.A,I.M., Lcmstra, PJ., Kanamoto, T. and Bastiaansen, C.W,M"

Polymer 1991, 32, 34

6. Nikitin, V.N., Volkova, L.A, Mikhailova, N.Y. and Baklagina, Iu.G.,

Vysokomol. soedin. "1959, 1, 406

7. Morera, D., Ciampelli, F. and Mantica, E., Adv. Malec. Spectrose., Prot. Int.

Meet. 1962, 2, 898

8. Moller, M., Makrornol. Chern. Rapid Comm, 1988,60/61, 107

9. Deckmann, H., PhD. Thesis, Albert-Ludwigs-Univerisiit, Freiburg, Germany,

1991, Chapter 4

10. Smith, P" Lemstra, P.l., Kalb, B. and Pennings, AJ., Polym.Bull, 1979, 1, 733

11, Smith, P, and Lemstra, P.l.,l. Mater. sci 1980, 1S, 505

12. Anadakumaran, 1<., Roy. S.K. and Manley, R.SJ., Macromolecules 1988,21,

1746

13. Termonia, Y., Meakin, P. and Smith, P., Macromolecules 1985, 18,2246

14. Nakamae, K., private communication

15. Bastiaansen, C.W.M" Ph.D. Thesis, Eindhoven University of Technology, The

Netherlands, 1991

Electron Beam i"adiaJion of Drawn UHMW·PE/T-I,4-PB Blends

Chapter 5

Electron Beam Irradiation of Drawn

UIIMW-PE/T-l,4-PB Blends

5.1 Introduction

79

Electron beam (EB) irradiation of oriented UHMW-PE tapes results in crosslin­

king~ hut also in chain-scissioning [1,2]. With increasing draw ratio, I.e. degree of

orientation and chain extension, the ratio of crosslinking over scission decreases

[3]. As a consequence EB"irradiation of highly oriented PE structures, results in

a decrease in properties like tensile strength and creep rate [4,5,6].

In the case of oriented UHMW-PE/t-l,4-PB blends, some conditions seem to be

fulfilled to obtain a more favourable ratio of crosslinking over chain scission: The

presence of t-1,4-PB with a high G-value for crosslinking [7] and a high segmental

mobility in the hexagonal phase, which can be considered as a macromolecular

prorad. Furthermore, in oriented tapes, t-1,4-PB is dispersed in the UHMW-PE

matrix on a very fine scale, as is shown in chapter 4. Though the two components

are not mixed on a molecular level, the t-l,4-PB domains are that small that they

can not be detected by DSC, TEM and W AXD. Only by spectroscopic techniques,

like NMR and IR, the presence of t-lA-PB domains in oriented blends can be

demons trated.

80 Chapter 5

In this chapter, the effects of EB irradiation on the properties of the oriented

blended systems are studied. For that purpose the influence of irradiation dose and

irradiation temperature, as well as the t-l,4-PB content were varied.

5.2 Experimental

5.2.1 Sample Preparation

The UHMW-PE used in this study was Hostalen Gur-412 supplied by Hoechst

Ruhrchemie (1\\..= 1,500 kg/mole). The t-1A-PB was synthesized in our laboratory

(M.,.=60 kg/mole, vinyl-content 0.9-1.2%). Experiments were performed with

UHMW-PE/t-1,4-PB blends containing 1, 5, 10 or 20 wt.% t-l,4-PB and 100%

UHMW-PE samples were used as a reference. The blends were prepared via

solution-mixing in xylene (1.5 % (w/v) polymer in xylene) and subsequent sOlution­

crystallization according to the method described in detail in chapter 3. The

resulting films were extracted with n-hexane in an ultra-sonic bath to remove the

stabilizer which was added to prevent crosslinking and degradation during

diS50lution. The films were cut into tapes (8 nun x 25 nun) which were drawn

manually to a draw ratio of 40 on a hot shoe at 120 °C using inkmarks to determine

the displacement. A draw ratio of 40 was chosen, because this draw ratio could

easily be obtained for all samples.

To study the off-axis properties of the drawn blends, films of UHMW-PE and of a

hlend containing 10 wt.% t-1,4-PB were drawn at 120 °C, to a draw ratio of about

25. Due to the large width of these films (80 nun), 25 was the maximum draw ratio

which resulted in more or less homogeneous tapes. After EB irradiation, the tapes

were cut perpendicular to the drawing direction into 5mall strips (2 nun x 50 mm)

and prepared for tensile testing as described in section 5.2.5.

Electron Beam Irradiation of Drawn l!!:£A!.W.PEfT-1,4-PB Blends 81

5.2.2 Electron Beam Irradiation

Irradiation was performed with a 3 MeV 'Van de Graaf acccierator at the Interuni­

versitair Reactor lnstituut, Delft. The tapes were mounted on aluminum plates and

placed in an aluminum heating box with a window inert for the electron beam. The

box was flushed with nitrogen gas before and during irradiation to prevent oxidative

degradation (chain scission). The blended tapes and reference PE tapes were

irradiated at 30°C or at 100 ~C (above the solid-solid phase transition of t-1,4-PB

and below the melting temperature of PE and H,4-PB) with doses of 20, 60, 100

and 150 kGy, respectively.

Some tapes were kept in the heating box at 30°C or 100 °C without irradiation for

the same time as the irradiated samples, as a reference for the influence of

temperature, i.e. relaxation, on the properties of the tapes.

Immediately after irradiation, the samples were placed in a dosed box flushed with

nitrogen. The samples were kept in this nitrogen atmosphere for 4-5 days to ensure

the decay of trapped radicals in the absence of oxygen, again to prevent oxidative

degradation. This procedure was used since mechanical properties like creep rate

and tensile strength are rather sensitive towards chain scission.

5.2.3 Gel Fradion and Swelling Ratio

To evaluate the crosslinking of the samples, gel fraction and swelling ratio of the

irradiated materials were determined via the procedure given below. Firstly, the sol

fraction of the irradiated samples was removed by soxhlet extraction with xylene for

at least 48 hours. The residual gel was kept in xylene at 120°C for 5 hours to reach

its equilibrium degree of swelling. After 5 hours, the swollen gel was transferred

into a stoppered flask and weighed, giving the weight of the swollen gel (W~)_

Subsequently, the swollen sample was dried under vacuum at 50 ~C and weighed

again (Wd).

82 Cha.pter 5 ------------------~----------------------------------

The gel fraction (fG) was calculated from the initial weight of the samples before

extraction (Wo) and the weight of the dried gel (Wd ) using equation 5.1.

(5.1)

The volume degree of swelling or swelling ratio (SR) of the samples was calculated

from Ws and W<j a.~!>uming additivity of volume in the swollen gel, using equation

5.2

(5.2)

in which Pp and Ps represent the densities of the polymer and the solvent, with

values at 120 °C of 0.908 gcm-] and 0.768 gcm-3, respectively. The density of the

polymer is not corrected for the presence of t-l,4-PB in the gel.

The gel fraction is a measure for the amount of material involved in the network

formed and the swelling ratio is a measure for the density of the network. The

values reported are averaged over 3 experiments.

5.2.4 Maximum Draw Ratio

After BB irradiation of drawn tapes (lJ '" 40) containing 10 wt. % t-l,4-PB, the tapes

were drawn further at 120°C to failure (Az). This latter draw ratio, (A z), is

dependent on the crosslinking efficiency. The maximum obtainable draw ratio (lmax)

is defined as the product of (11) and (A2).

The values reported are averaged over 4 to 5 experiments.

Elec;tron Beam Irradiation of Drawn UHMW-PEjT-l,4-PB Blends 83

5.2.5 Tensile Testing

Tensile testing was performed at room temperature on a frank 81565 tensile tester

equipped with an e~tensometer. To avoid slippage in the damps, the tapes were,

on botb ends, glued between cardboard tabs. The strain rate adopted was 10-3 S-1

with an initial length between clamps of 150 mm. The small strips used to test the

off-axis properties possessed an initial length between clamps of 25 mIn. The values

reported for Young's modulus and tensile strength are averaged over 5 experiments.

5.2.6 Creep Measurements

Creep measurements were performed at 30 <)C, with irradiated blended tapes con­

taining 10 wt.% t-l,4-PB and with pure polyethylene tapes, using a Frank 81565

tensile tester equipped with a thennostaticaIly controlled oven. The tensile tester

was adopted for dead-loading creep measurements. A constant stress of OA GPa

was applied. Elongation was measured as a function of time using an extenso meter.

The initial length between clamps (10) amounted to 100 nun. Plateau creep rates

were determined from the constant slope of the strain-time curves; the values

reported are averages over 2 to 3 experiments.

5.2,7 Differential Scanning Calorimetry (DSC)

Thennograros were recorded using a Perkin-Elmer DSC-7 differential scanning

calorimeter. The experimental parameters were described before in chapter 3.

84 Chapfer 5

5.2.8 Constrained Heating

To establish whether a 'macroscopically' homogeneous network has been formed

by EB irradiation, the tapes were subjected to a constrained heating test. The

irradiated tapes were clamped at both ends in a brass frame, thus keeping the

samples at a constant length of 150 mm. This frame was placed in an oven at 200

DC for 10-30 seconds. If no homogeneous network is formed by EB irradiation the

tapes will melt, break and shrink. However, if a network is present, the tapes will

remain intact. The residual tensile and thermal properties of the latter tapes were

determined according to the methods described above in section 5.2.5 and 5.2.7,

respective ly.

5.3 Results and Discussion

5.3.1 Crosslinking Efficiency

in figure 5.1 the gel fraction of the tapes is given as a function of both t-l,4·PB

content and irradiation dose, after irradiation at 30 OC (figure S.la) and at 100 DC

(figure 5.1b). Please note that the abscissa are not linear.

Compari~on of figures 5.1a and S.lb shows that the gel fractions obtained by

irradiation at 100 DC are, a5 expected, higher than after irradiation at 30 DC. All

samples irradiaLed at 100 DC have formed a network already at the lowest dose of

20 kGy, indicating that the gel-point dose (the dose at which a gel is formed) is <

20 kGy. This in contrast to the samples irradiated at 30 DC, which show only

network formation at 20 kGy if they contain more than 5 wt. % t·l,4.PB. In the

literature, the increased crosslinking efficiency upon irradiation at high temperatures

is contributed to an increased chain mobility and higher reaction rates [6,8].

Elec;tron Beam Irradiation of Drawn UHMW·PE/T·J,4.PB Blends 85

100

i ~ 80

~ "'"' 60 !5 ()

~ 40

20

0

100

~ 80

~ ... 60 !5 ()

3 40

20

0

Figure 5.1

0 1 5 10 20 t-1.4-PB content <wt.%)

(.:77J 20 kG-y

W¥J 60 kGy 0 5 10 20

t-1.4-PB content (wt.%) ., 100 kGy

_ 150 kGy

Gel content of blend tapes as a function of t-l,4-PB content and

irradiation dose. Irradiation peiformed at 30 DC (a) and at 100 DC (b)

86 Chapter 5

The influence of irradiation temperature is less pronounced for samples containing

more t -1,4-PB, hecause in these samples the crosslinking efficiency is already high

at room temperature. This is partly due to the presence of more unsaturated bonds

and partly to the fact that the t·l,4·PB domains in the drawn tapes are, even at

room temperature, largely in the mobile hexagonal and/or amorphous phase (see

section 4.3).

The density of the network is inversely proportional to the swelling ratio. The

swelling ratios of the soxblet residues discussed above, are given in figure 5.2 as a

function of both t-l,4-PB content and irradiation dose. Please note that, because of

the inverse relation between gel-fraction and swelling ratio, the axes of the

irradiation dose in figure 5.1 and 5.2 are inverted.

The higher crosslinking efficiency upon irradiation at 100 QC is illustrated in these

experiments as well, albeit less clear, since the swelling ratio of the samples

irradiated at 100 QC are lower than those irradiated at 30 "c. Both figures show the

dear influence of irradiation dose, i.c. higher doses resulting in lower swelling

ratios.

An apparent maximum in swelling ratio can be observed for blends containing

5 wi. % t-l,4-PB, suggesting a lower crosslinking efficiency in these blends. The same

effect, though less pronounced, can be observed in figure 5.1 for the gel fraction

which seems to show a corresponding minimum around 5 wt.% t-l,4-PB. No expla­

nation for this phenomenon has been found, yet, but the same effect, a maximum

or minimum around 5 wt.% t-l,4-PB, has been observed for other properties (see

e.g. the melting temperature after irradiation in figure 5.7b).

Another method to study the cross linking efficiency is to determine the maximum

drawability after EB irradiation. In figure 5.3 the results of such an experiment with

a sample containing 10 wt.% t·l,4·PB. are shown for irradiation at 30 ~C and at

100 ~c.

Electron Beam Irradiation of Drawn UHMW-PEjT-l,4-PB Blends 87

120

I 100

0 .;=,

80 !Il ....

~ 60 m 3: (J)

40

20

0

120

I 100

.Q - 80 ~

,~ 60

~ (J)

40

20

0

Figure 5.2

0 1 5 10 20 t-1.4-PB content (wt.%}

0 5 10 20 t-1A-PB content (wt.%) .. 100 kGy

_ HIO I<.Gy

Swelling ratio of blend tapes as a [unction of t-l,4-PB content and

irradiation dose. Irradiation performed at 30 "C (a) and at 100°C (b)

88 Chapter 5 . __ .",'\----

200,-----------------------.

I 150 0

\ \ :;:; \ (lj

<,. \ \

~ \:-ro 100 T t '\ '\

'" t::. § '" T ""- ",,- ,.. l!,. E .L

X ~"-- -g 50 .................. 2, 5 ----- -""C..._--..

1 L-______ ~~ ______ ~ ________ ~ ____ ----~

o 50 100 150 200

irradiation dose (kGy)

Figure 5.3 Maximum draw ratio (A'1t! '" A I.x Aj A,;;;;;; 40)) of a blend tape contain"

ing 10 wt. % t-l,4-PB as a function of irradiation dose (irradiated at

30 ~C (11) or 100 ¢C (0))

These results confirm the trends observed in swelling ratio and gel fraction curves,

Le_ a decrease in maximum drawability with increasing irradiation dose, and thus

an increase in crosslinking efficiency. 1bis effect is slightly more pronounced after

irradiation at 100 "C than at 30 "c.

5.3.2 Mechanical Properties

The influence of EB irradiation on the Young's modulus of the tapes is given in

figures 5.4a and 5.4b for tapes irradiated at 30 <:>C and at 100 "C, respectively. On

an average, the Young's moduli of tapes irradiated at 100 "C are slightly lowe\" than

those of tapes irradiated at 30°C, Since also the reference tapes, which were kept

Electron Beam Irradiation of Drawn UHMW-P£jT-I/4.PB Blends 89

at 100 ~C without irradiation, show this effect, this slight decrease is related to

relaxation processes at 100 °C resulting in a decrease in chain orientation and/or -

extension. Except for this temperature effect, the differences in Young's moduli of

unirradiated tapes, tapes irradiated at 30 Q and tapes irradiated at 100°C are not

very large and usually within the accuracy of the measurement. Therefore, it can be

concluded that the Young's modulus of these tapes is not significantly effected by

EB irradiation which is consistent with observations reported in literature for pure

polyethylene fibres [4,5].

The tensile strength of the tapes is more influenced by EB irradiation (see figures

5.5a and 5.5b) than the modulus, at least in the case of low t-l,4-PB contents_ The

tensile strength of tapes containing 0 to 5 wt. % t-l,4-PB decreases with increasing

irradiation dose. The tensile strength of tapes containing more than 10 wt. % t-1,4-

PB however, hardly changes upon irradiation. These effects are the same for

irradiations perfonned at 30 QC and at 100 ac. After irradiation, all blends possess practically the same tensile strength (0.8 to 0.9

GPa) regardless the initial value.

In the past, similar decreases in tensile strength have been observed upon

irradiation of drawn UHMW-PE samples [4,5]. In these samples the decrease in

tensile strength upon irradiation has been attributed to main-chain scissioning,

causing an increase in the amount of defects and loose chain ends which has a

diminishing effect on tensile strength.

The fact that the tensile strength of blends containing more than 10 wt. % t -1,4-PB,

does not decrease upon irradiation can be explained by a lower degree of chain

scission in these blends due to the presence of t-l,4-PB.

90

fa 0.. ~ I/) :J "5 "0

~ If)

.~ >-

I.f) ;;l

"5 "0 o E

Figure 5.4

70

60

50

40

30

20

10

a 0

0 5 10 15 20 25

t- 1.4-PB content (wt.%)

70 ~~-------------------------------,

60

40

D

10

'. , ··· .. '6

V

b

Chapter 5

o ~ ____ ~ ______ -L ______ ~ ______ L-____ ~

o 5 10 15 20 25

t-1,4-PB content (wt.%)

Young's modulus of blended tapes (L\' a function of t-l,4-PB content

and irradiation dose; 0 (0), 20 (6.), 60 (0) and 100 kGy (V).

Inudiation pe!formed at 30 QC (a) and at 100°C (b)

Electron Beam Irradiation of Drawn UHMW-PEjT-1,4-PB Blends 91

ro ~ £. ~

~ "-... III

~ "@ 2

Figure 5.5

1.60 .---------------------,

a 0.00 L.....-__ -----L_~_ ......... _-___._L ___ ...l.._ __ ____'

o 5 10 15 20 25

t-1,4-PB content (wt.%)

1.60 ....--------------~-----,

1.20 \. 1J:r--- " ~

G---7'....zi.:.....-:-_:.~ .... . " " "

"'1;7"

0.80

0.40

b 0.00 L-__ ----L __ ~...L....._~__L ___ ....L.... __ ___..J

o 5 10 15 20 25

t- 1,4-PB content (wt.%)

Tensile strength of blended tapes as afimction oft-l,4-PB content and

irradiation dose; 0 (0), 20 (11), 60 (0) and 100 kGy (V). 'rradiation

peiformed at 30 °C (a) and at 100 °C (b)

92 Chapter 5 ._----------_ ..... , ... _-

Figures 5.6a and 5.6b show the effect of irradiation dose on the creep rate of a

blend containing 10 wt.% t·l,4.PB upon irradiation at 30 °C and 100 °C re~pcctively.

As a reference the creep rates of comparable 100 % UHMW-PE samples are

shown in the ~ame figures.

These resu1t$ clearly show an increase of the creep rate of pure polyethylene tapes

with increasing irradiation dose. Similar results were Obtained in the past via EB

irradiation of gel"spun fibres [9,lOj. The increase in creep rate is generally

attributed to an increase in the degree of chain scissioning with increasing dose, as

in case of the tensile strength. This effect is leS5 pronounced upon irradiation at

100 °C because the higher chain mobility at this temperature increases the degree

of crOSSlinking at the expense of the degree of chain scissioning.

The addition of 10 wt. % t-1.4-PB has no significant influence on the creep rate as

far as unirradiated tapes are concerned. However, upon irradiation the presence of

t-l,4-PB has a considerable effect on the creep rates. The drastic increase in creep

rate encountered in PE tapes upon irradiation is prevented by the presence of t-1,4-

PB, resulting in creep rates which are only slightly higher than in unirradiated

blends. Consequen tly, the creep rate of a blend containing 10 wt. % t· 1 ,4-PB is only

slightly affected by EB irradiation just like the tensile strength. Since both

properties (creep rate and tensile strength) are influenced by chain scissioning, and

hlends containing less t-l,4-PB do show decreases in tensile strength upon

irradiation, these results point to an ability of t·1,4.PB to decrease the amount of

chain :sci~~ion in the blends upon irradiation. A:s far as chain-scissioning is

concerned, t-) ,4.PB appears to work as a stabilizer. Despite the fact that the

presence of t-l,4-PB prevents a. drastic increase of creep rate upon irradiation, it

does not irnprove the creep re~istance of the blends. Since gel fraction and swelling

ratio measurements have demonstrated the presence of a network in the:se blends,

the question arises how homogeneous this obtained network is and why it does not

interfere with the creep mechanism. These que:stions will be addressed in chapter 6.

Electron Beam Irradiation of Drawn UHMW-PE/l'-1,4-PB Blends 93

50

40 ~

'i (f) ,. 0 30 ,....

GJ ... ro ....

20 ar GJ tI

10

0

50

40 ~ , (f)

'!' 0 30 .,--

GJ .,.... l'!l '- 20 g ~ (J

10

o

Figure 5.6

IIlIIliifIbl~

C:·:·::·:·:I PE

a ..

0 20 60 irradiation dose (kGy)

o 20 60 irradiation dose (kGy)

Creep rate oj a 100 % PE tape and a blended tape containing

10 wt.% t-1,4-PB as a junction oj irradiation dose. Irradiation

peiformed at 30 °C (a) and at 100 °c

94 Chapter 5 -------_ ........ ,." ... " ..... __ ... _--------------- .. "" ...... _-

5.3.3 Melting Behaviour and Constrained Heating

The influence of BB irradiation on the melting temperatures of the tapes is shown

in figures 5.7a <lnd 5.7b. The melting temperature of the tapes clearly decre,lses

with increasing irradiation d05e, both upon irradiation at 30 °C and upon irradiation

at 100°C. This decrease in melting temperature is due to the introduction of a

number of defect::; (e.g. chain ends, crosslinks and branch points) in the structure

by 8H irradiation [11].

The inHuence of t-'1,4-PB content in this case is less obvious. In general, it seems

that in blends with a higher t-l,4-PB content the melting temperature is decreased

after irradiation. However, a maximum can be observed for blends containing 5

wt.% t-l,4.PB, especially for blends irradiated at 100°C (see figure 5.Th), for whieh

no explanation has been found yet (see also figures 5.1 and 5.2).

The crystallinity of the samples is not influenced significantly, all values vary around

70 %. Considering the maximum dose used (only 100 kGy) no changes in

crystallinity would be expected [8,12].

A" was shown in chapter 4, figure 4.4, the solid-solid phase transition of t-l,4.PB

disappears from the DSC thermograms upon drawing. This was attributed partly,

to the transition of monoclInic into hexagonal material upon drawing and partly, to

the uecrea:-;e in the dimensions of the t-l,4-PB domain:;. In the second DSC run

however, after melting and recrystalli7,ation, this transition reappears in the DSC

curves. Thi:; latter effect does not occur after BB irradiation. Apparently, irradiation

of the oriented hlends leads to such changes in the molecular structure of t-l,4-PB

(like branching, crosslinking and loss of unsaturation) that even after melting and

recrystallization, the solid-solid phase transition can not occur.

Electron Beam Irradiation of Drawn UHMW-PE/l'-1,4-PB Blends 95

2 145

<l> 144

~ .- 143 m iii g. 142 <l> -I-'

141

~ .... 140 1) E 139

138

137

2 145

<Ii 144 j -I-' 143 ro

~ 142 Q>

...." 141

~ .;::; 140 Q)

E 139

138

137

Figure 5.7

0 5 10 t-1,4-PB content (wt.%)

0 1 5 10

20

20

~ 100 kGy

j:.. ... ;;:.;.,': ... ;;J 60 kG:,;

c:::J 20 kGy

t-1,4-PB content (wt.%) [::J 0 kGy

Melting temperature of blended tapes as a function oj t-l,4-PB content

and irradiation dose. Irradiation perj'onned at 30°C (a) and at 100 °c (b)

96 Chapter 5

Constrained heating of the irradiated samples according to the method described

in ~cction 5.2.8, demonstrates a clear advantage of the modification of PE fibres

with t-l,4-PB. Only irradiated blended tapes remain intact in this test. All un­

irradiated samples and irradiated samples without t-l,4-PB, melt completely and

break. In the past, Hikmet et al. [13] reported similar results for UHMW-PE

fibres prepared from gel films which were irradiated prior to drawing. However, due

to this procedure the drawability is limited.

The constrained heating experiment shows that the network present in the

irradiated blend samples is homogeneous enough to yield a macroscopica.lly

coherent structure.

In chapter 6 the topology of the network obtained will be discussed in more detail.

The residual properties of the samples which do remain intact during constrained

heating might provide some mOre insight into the structure of the network.

5.3.4 Residual Properties after Constrained Heating

Tables 5.1 and 5.2 show the residual Young's moduli and tensile strengths after con­

strained heating, of blends irradiated at 30 °c and at 100 0c. The modulus of the irradiated blend samples strongly decreases upon constrained

heating from values between 50 and 30 GPa (depending on the t-l,4-PB content)

before constrained heating to values between 10 and 15 GPa after constrained

heating. The decrease in modulus is larger for blends with low t-l,4-PB contents and

for irradiations performed at 30 0c. The irradiation dose also has a slight effect,

higher doses resulting in smaller decreases- The tensile strength of the blends shows

a similar behaviour as the modulus upon constrained heating, only the relative

decreases in tensile strength are smaller, as expected. In general, the modulus drops

around 70 % whereas the tensile strength decreases 30 % at most.

Electron Beam Irradiation of Drawn UHMW-PEjT-l,4-PB Blends 97

Table 5.1 Residual modulus and tensile strength of drawn blends after irradiation

at 30 °C and subsequent con..~trained heating at 200 °C for 10 seconds

jrrad. dose t-l,4-PB content E mOdulus tensile strength

(kGy) (wt.%) (GPa) (GPa)

20 1 12 0.8

5 15 0.8

10 11 0.7

20 12 0.6

60 1 14 0.8

5 9 0.7

10 15 0.7

20 13 0.5

100 1 17 0.7

5 13 0.6

10 15 0.7

20 14 .0.5

These results show that blends with high t-l,4-PB contents, irradiated with high

doses at toO OCt are best capable of retaining the orientation and extension of the

chains, upon constrained heating. This implies that the network is most effective in

these cases. Unfortunately these are the blends with the least impressive tensile

properties after EB irradiation.

98 CJlapt(~r 5 -------_ ........... _--------- -------

Table 5.2 Residual modulus and temile strength of drawn blend~ atia imu:iiation

at 100°C and subsequent constrained heating at 200 OCfor J() seconds

in-au. dose t-l,4-PB content E modulus ten:;ile strength

(kGy) (wt.%) (GPa) (GPa)

20 14 1_0

5 12 0.8

10 10 0.8

20 14 0.6

60 '[ 21 0-9

5 10 0_7

to 16 0-8

20 17 0.7

100 19 0_8

5 14 0.8

10 20 0.7

20 15 0-7

The drop in mechanical properties must be caused by a significant change in

morphology. Transmission electron micrographs of the blends after constrained

heating clearly show the presence of shish-kebab like structures (sec e_g_ figure

5.8a). Higher magnification reveals long crystalline cOres with tapered and

interlocking lamellar 'branches' (figure 5.8b).

Electron Beam Irradiation of Drawn UHMW-PE/T-l,4-PB Blends 99

Figure 5.8 TEM micrographs of an irradiated blend tape after constrained heating

at 200 °C

The occurrence of this shish-kebab structure is due to melting of chain ends and

other parts of the network which are not under stress during constrained heating.

These molten chains recrystallize as chain-folded crystals on the loaded extended

cores upon cooling after constrained heating.

The existence of a shish-kebab structure also appears form the DSC thermograms

of the samples after constrained heating. These thermograrns show that the peak

melting temperature (due to extended-chain crystals) only slightly shifts towards

lower values but that a shoulder appears on the left hand side of the peak around

100 Chapter 5 ----------------------------------------------------------

135 ~C which is due to the introduction of chain-folded lamellae (compare rderence

[14]). An example of this effect is given in figure 5.9.

Figure 5.9

--'~~.~--------------,

100

a

b

150 ----Temperature lOC)

DSC thennogram of a blend tape containing 5 wt. % t-l,4-PB,

irradiated at 100 QC with 20 kGy, 0) before constrained heating and b)

after constrained heating

The overall crystallinity of the samples after constrained heating is on average about

10 % lower than before constrained melting which can be readily explained by the

transition of a part of the chain-extended crystals into chain-folded crystals.

The observations concerning the residual properties and morphology of the

irradiated blends after constrained heating, clearly demonstrate that only part of the

blend is involved in a continuous, stretched and oriented network. Furthermore, the

fact, that the residual properties obtained can already be achieved when only 1

wt.% t-I,4-PB is present in the blends, indicates that the the continuity of the

network exists through a relatively small amount of t-l,4-P13. Adding more t-l,4-PB

results in larger t·1,4·PB domains which, in view of the residual properties, do not

Electron Beam I"adiation of Drawn UHMW.PEjT-l,4-PB Blends 101

add much to the network effectivity in retaining orientation and enxtension of the

PE chains.

5.3.5 Transverse Properties

Tables 5.3 and 5.4 give the modulus and tensile strength in the direction perpen­

dicular to the drawing direction of about x 25 drawn tapes of pure polyethylene and

of a blend containing 10 wt.% t·l,4·PB. These results show that the network

introduced by the presence of t-l,4-PB during irradiation has no influence on the

transverse properties.

Tobie 5.3 Transverse modulus and tensile strength of PE-tapes (k ... 25) as a

junction of imuliation dose (in-adiated at 30 "C)

irrad. dose E-modulus tensile strength elong. to break

(kGy) (GPa) (MPa) (%)

0 0.6 9-9 15

20 0.7 8.2 1.7

50 0.9 9.5 1.4

100 0.5 4.0 1.1

102 Chapter 5 __ •• w- . _________ _

Table 5.4 Tranwerse modulus and tensile strength of FE/PS·tapes (..t ;;;;;25, /-1,4-

PE content ::::: 10 wt. %) as a function of ifmdiatioll dose (irradiated

aJ 30 ~C)

irrad. dose E·modulus tensile strength elong. to break

(kGy) (GPa) (MPa) (%)

0 0.9 10.4 2.0

20 0.6 8.0 1.5

50 0.8 7.7 1.3

100

Considering the observations described above, this should be expected. Since not

all material is involved in the network, the transversal properties are still completely

determined by the low intermolecular Van der Waals interactions.

5.4 Conclusions

Electron beam irradiation of drawn UHMW-PE!t·l,4-PB blends results in the

introduction of a netvvork. The amount of material involved in this network and the

network density increase with increasing irradiation dose and with increasing

irradiation temperature.

The modulus of the tapes is not significantly influenced by ES irradiation. The

tensile strength of pure polyethylene lapes and blended tapes with low t·l,4·PB

contents, decreases upon irradiation whereas the blended tapes with higher t·l,4-PB

contents (> 10 wt.%) retain their tenacity upOn irradiation.

The presence of t-l,4.PB appears to reduce the amount of chain scission and to

lower the gel point dose of the drawn blends compared to pure UHMW-PE

Electron Beam Irradiation of Drawn UHMW-PE/T-l,4-PB Blends 103

samples. In this way it becomes possible to introduce a macroscopically

homogeneous network in the blend tapes.

This network can sustain constrained heating to > 200 DC, resulting in shish-kebab

structures, whereas irradiated pure UHMW-PE samples melt and break under these

conditions. The introduced networks however, are not capable to improve the creep

resistance of the fibres nO.- tbe transversal properties.

5.5 References

1. Charlesby, A, 'Atomic Radiation and Polymers', Pergamon Press, Oxford, 1960

2. Dole, M., 'The radiation Chemistry of Macromolecules', Academic Press, New

York, 1972

3. van A~rle, N.AJ_M_, Crevecoeur, G. and Lemstra, P.I., Polym. Comm. 1988,

29, 128

4. de Boer, 1. and Pennings. AJ., Polym. BulL 1981, S, 317

5. de Boer, 1. and Pennings, AJ., Coli. Polym- Sci. 1983,261, 750

6. Dijkstra, D.l. and Pennings, AJ., Polym. BulL 1987, 17,507

7. van Gisbergen, J.G.M., PhD- Thesis, Eindhoven University of Technology, The

Netherlands, 1991, Chapter 2

8. O'Donnell, I.H., 'Effects of Radiation on High.Technology Polymers', ACS

Symp. Ser. 1989,381, Chapter 2

9. de Boer, J. and Pennings, AI, Po/ym. Bull. 1988, 19, 73

10. Klein, P-G-, Woods, D.W. and Ward, 1M., 1. Polym. Sci. Polym. Phys. Ed.

1987, 25, 1359

11. Mukherjee, AK., Gupta, B.D. and Sharma, P.K.,lMS-Rev- Macromol. Chern.

Phys. 1986, C26, 415

12. Keller, A, 'Developments in Crystalline Polymers-l', ed. by Bassett, D_C.,

104 Chapter 5 -.~,,,, .. ----

Applied Science Publishers Ltd., London, 1982, Chapter2

13. Hikmet, R, Lem$tra, PJ. and Keller, A, Col/. Polym. Sci. 1987,265, 185

14. de Boer, J, van Hutten, P.F. and Pennings, A.J., 1 Marer. Sci. 1984, 19, 428

'i!!:::::ork Topolqs:

Chapter 6

Network Topology

6.1 Crosslinking in Retrospect

105

In chapter 5 it was shown that electron beam (EB) irradiation of drawn UHMW­

PE/t-l,4-PB blends results in the formation of a network. The gel fraction and the

network density increase with increasing dose and temperature of jrradjatjon. The

presence of t-l,4-PB lowers the dose necessary to obtain a gel (gel point dose) of

the system and reduces the amount of chain scission, which appears from a smaller

decrease in tensile strength upon irradiation of blends compared to pure

polyethylene.

The presence of a network in the drawn t-l,4-PB/UHMW-PE tapes is not only

apparent from the gel fraction measurements but is also demonstrated by

constrained heating experiments. Irradiated, drawn t-l,4-PB/UHMW-PE blends

remain dimensionally intact upon constrained heating to approximately 200 DC,

whereas comparable irradiated pure UHMW-PE tapes fail under these conditions.

106 Chapter 6 ~--------------- ...... -,,--.-----

After constrained heating, shish.kebab structures are formed with a residual

modulus and tensile strength, which are, at most, respectively 70 % and 30 % lower

than before constrained heating of the irradiated samples.

Though the network obtained jmproves the thermal resistance of the oriented

blends it is not capable of improving the creep resistance and the transverse

properties of these samples.

Consequently, the paramount question is where the crOSSlinks are located in the

blended drawn system, i.e. what the network topology is, which is directly related

to the question how the t-l,4-PB is dispersed in the system.

In chapter 3 it was established that prior to drawing, the t-I,4-PB domain:; are

located between stacked polyethylene crystals, the dimensions of the t" 1,4-PB

domains are dependent on the t-l,4-PB content in the blend. Upon drawing, the

polylmtadiene domains are elongated via plastic deformation imposed by the

polyethylene matrix and finely dispersed within the system. Due to this deformation,

the dimensions of the t-l,4-PB domains decrease to such an extent, that in case of

the blends with lOw t-lA-PB contents « 20 wt.%) they can not be detected by

DSC, WAXD and TEM at high draw ratios.

Based on these observations one is tempted to describe the structure of the drawn

blends by a simple two-phase model comprising elongated t-IA-PB domains finely

dispersed in between the chain-extended PE fibrils of the matrix.

However, such a simple two phase model can not explain the observed network

continuity. Therefore, considering the macroscopic homogeneity of the network

obtained, even when only I 'Wt.% t-I,4-PB is added, we assume that some t-I,4·PB

chains must be dispersed on an even finer scale within the drawn blends, in order

to account for the continuity of the network throughout the blend tapes.

Since the chain"extended PE fibrils are not 100 % crystalline but contain numerous

defects related to trapped entanglements, chain ends, tie-molecules etc., which are

present throughout the whole tape, it is assumed that a minor fraction of the

polybutadiene chains is trapped within these amorphous regions. Upon FH

Network Topology 107

irradiation, these trapped chains provide for continuity within the blended systems

via the formation of crossIinks.

In figure 6.1 a schematic drawing is presented of the structure of the crosslinked

dravm blends.

Figure 6.1

(=) defect ~ODe • ~-l.+polybutadi= molecule

elongated, interfibrillar. t-l.+PB domain

polyethylene fibril

Schematic representation of the strncture of CTOsslinkec4 drawn UHMW­

PE/t-1,4-PB blends

In chapter 5 it was shown that the residual modulus of the blended tapes, irradiated

with a certain dose, does not vary with t-1,4-PB content (see tables 5.1 and 5.2).

This indicates that already at 1 wt.% t-l,4-PB a maximum effective network is

obtained. Adding more t·l,4-PB to the blends does not increase the amount of

108 Chapter 6

effective crosslinks, i.e. the number of load bearing chains, because the surplus of

t-lA-PB agglomerates in dispersed t-l,4-PB domains, not further expanding the

netork.

Increasing the irradiation dose and/or -temperature of a blend with a given t-l,4·PB

content, slightly increases the effects of crosslinking as is reflected by a slightly

higher residual modulus. However, the amount of chain scission increases more

pronounced, resulting in an undesirable decrease in residual tensile strength.

1.2 Crosslin king in Prospect

The creep resistance of the oriented systems might be improved if the number of

crosslinks could be increased, and if the crosslinks would be distributed more

homogeneollsly throughout the samples. Since the creep mechanism is assumed to

be related to slippage of chains through the PE crystal lattice [1,2,3,4,5), it is

probably necessary to introduce cross links in the crystalline regions to obtain an

increase in creep resistance at alL

The density of the netW(lrk obtained at a certain irradiation dose, might be

increased by increasing the amount of t-l,4-PB in the disordered domains_ This can

possibly be accomplished by the use of a solvent which is less discriminative to

UHMW -PE and t-1,4-PB than xylene, as far as crystallization is concerned, since in

chapter 3 it was shown that upon blending of UHMW-PE and t.l,4-PB in solution,

crystallization induced phase separation (L-S) occurS.

Another possibility is to increase the compatibility of UHMW-PE and t-l,4-PB via

partial hydrogenation of t-l,4-PB. In that case, th~ amount of t-l,4-PB trapped

within the amorphous domains might be higher. For this purpose also the use of

other (partially hydrogenated) polyalkenamers, like trans-polyoctenylene, could he

contemplated.

Network Topology 109

Further hydrogenation may eventually lead to the incorporation of some

unsaturated C-C bonds in the crystalline regions which in principle could enhance

the homogeneity of crosslinking. The question remains however, whether

crosslinking will occur at all within the crystals, since the chains are too far apart

in the orthorhombic crystal lattice of PE, to crosslink. Since segmental mobility of

the chains in the hexagonal lattice is much higher, a two-stage irradiation process

could be contemplated in which the first step introduces a network via iJTadiation

below the melting temperature. The presence of this network will facilitate a second

irradiation step at temperatures above the orthorhombic-hexagonal transition which

maybe permits crosslinking in the (hexagonal) crystalline regions_

Extrapolation of the results described in this thesis to the production of fibres via

the solution-spinning technique, hi not possible without further preface, in view of

different processing parameters involved [6].

6.3 References

1. Wilding, MA and Ward, I.M., Polymer 1978, 19, 969

2. Wilding, M.A. and Ward, I.M., Polymer 1981, 22, 870

3. Wilding, M.A and Ward, I.M., Plast. RUbb. Proc. Appl. 1981, 1, 167

4. Ward, I.M. and Wilding, M.A.,1 Polym. Sci. Polym. Phys. Ed. 1984,22, .561

5. Govaert, LE., Ph.D. Thesis, Eindhoven University of Technology, The

Netherlands, 1990

6. DSM Stamicarbon, NL 9.001.069, 1990

Morphology of Nasn:nt U!:!MW-PE 111

Appendix·

Morphology of Nascent UHMW-PE: Chain-Extended vs. Chain-Folded Crystals

A.I Introduction

Recently, the interest in the morphology of nascent reactor powders revived, due

to the observation by Smith et aI. [1,2,3], that Ultra-High MoleCUlar Weight

Polyethylene (UHMW-PE) reactor powders can be remarkably ductile, despite their

high molar mass. Smith et al. [1,2,3] showed that low temperature compacted films

of some nascent UHMW-PE reactor powders could easily be drawn into high­

modulus structures in contrast to compression moulded (melt-crystallized), but

chemically identical, films.

UHMW-PE powders with a high degree of drawability in the solid state could only

be obtained under specific polymerization conditions like low polymerization

temperatures, reduced catalyst activity and reduced mOnomer pressure [3]. In line

with the current view that solid-state drawability, at least in the case of poly­

ethylene, is limited by the presence of trapped entanglements [4J, these polymeri­

zation conditions promote disentangling as a consequence of 'free' chain growth.

* Reproduced in pan from: Y.M. T. Tervoorl-Engelen and P.J. Lemstra, Po/ym.

Comm. (1991), ,ll. 343

1"12 Appendix

In the past, the morphology of as-polymerized or nascent polyethylene has been

studied extensively [5,6,7,8,9]. Depending on the polymerization conditions,

morphologies ranging from globular to worm-like structures have been produced [5].

and on a sm'lller ~calc, shish-kebab-like structures could be observed occasionally

[7].

The melting temperatures reported for most of these nascent polyethylenes are in

the range of 138 to 143 OC [1,2,3,6,1O}, which is rather high compared to the

melting temperature of standard melt-crystallized polyethylene~ (133 - 135 0c). The

high melting temperatures of these nascent polyethylenes are generally attributed

to chain-extended crystallization during the polymeri:o.:ation process [2,6,7,9]. This

explanation can be supported by the absence of any clearly distinguishable Small­

Angle X-ray Diffraction maxima [6,9].

In this appendix, results will be presented concerning the morphology and melting

behaviour of some UHMW·PE reactor powders possessing the typical drawing

characteristics of highly disentangled PE-powders, as described by Rotzinger et at

[3}. These results will demonstrate that the commonly accepted relationship

between a high melting point (as measured by DSC) and the presence of (rather)

chain-extended crystals in PE samples can not be generalized.

A.2 Experimental

A.2.1 Materials

The nascent UHMW·PE reactor powders used for this study were polymerized with

a soluble vanadium (III) acetyl acetonate/diisobutyl alumjnium chloride catalyst at-

7 "C, according to the method described in detail by Rotzinger at a1. [3}.

Morphology of Nascent UHMW-PE 113

A.2.2 Differential Scanning Calorimetry

The thermal characteristics of the polymers were examined using a Perkin-Elmer

DSC-7 Differential Scanning Calorimeter with a standard heating rate of 1 0 ~C/ min.

The sample weight was approximately 5 mg. Indium was used for temperature and

heat of fusion calibration (Tm "" 156.6 ac. AHf '" 28.4 Jig).

The melting temperatures mentioned in this paper refer to the peak temperatures

in the thermograms unless stated otherwise.

A.2.3 Transmission Electron Microscopy

The polymer powder was embedded in an epoxy matrix, trimmed ready for

microtoming and subsequently treated during 16 hours with a ruthenium tetraoxide

solution prepared according to Montezinos et al. [11]. Finally, thin sections were

obtained at room temperature using a Reichert Ultracut E microtome. Transmission

Electron Microscopy was performed using a JEOL 2000 FX electron microscope

operating at 80 kV.

A.2.4 Average Crystal Sizes

The average crystal si2;es (ACS) of the reactor powders were obtained from X-ray

line-broadening data. Profiles of the (200) and (020) reflections were recorded using

a Philips PW1820 wide-angle diffractometer employed in the reflection mode. Ni­

filtered Cu I<,,; radiation from a Philips PW1731 generator operated at 50 kY and

40 rnA was employed. Instrumental broadening (b.J3jn• t) was corrected for using an

aluminium oxide powder. The half widths of the diffraction peaks were calculated,

assuming a Gaussian profile, with:

114 Appendix

(AI)

in which .1..Bwmpl~ is the broadening caused by the crystallites within the sample and

6.Bob~' the observed peak width at half-maximum intem;ity_

The average crystal sizes along the normals to the (hkl) crystal planes, Dhk1, were

calculated using the Scherrer equation:

0.9 ·1 D " .<

ItJrI .1..B • cose $ll"'pi. • ItJrI

(A.2)

in which 9 hk1 is the corresponding Bragg reflection angle and Ax is the wavelength

of the X-ray source (0.154 nrn).

A.2.5 Electron Beam Irradiation

Irradiation was performed using the electron beam of a Van de Graaff generator

at the Interuniversitair Reaktor Instituut (IRI), Delft. TIle samples were irradiated

in air at room temperature with a total dose of 250 kGy.

A.3 Results and DisCllssion

A representative sample of the nascent reactor powders described in the

Experimental section, showing the typical drawing characteristics as described by

Smith et al. [1,2,3], was selected. After compression moulding this powder below T m'

at 120 QC, the compacted film could be drawn to a draw ratio of 60 resulting in a

Young's modulus of 115 GPa.

Morphology of Nascent UHMW-PE 115

Figure A.1 shows a transmission electron micrograph of this selected polyethylene

reactor powder. This micrograph reveals that the crystals are not regularly stacked

which explains the absence of any Small Angle X-ray Diffraction intensity. Accor­

ding to the micrographs, the crystals possess a thickness of about 10 run and lateral

dimensions up to at most 100 run. The average lateral crystal sizes determined by

X-ray diffraction, based on (200) and (020) reflections, are 12.7 om and 11.6 nm,

respectively. Since the average lateral crystal sizes, determined by X-ray diffraction,

are usually quite inaccurate due to crystal imperfections, it is impossible to conclude

from these observations whether the crystals consist of chain-folded lamellae with

a thickness of 10-12 run and lateral dimensions of about 100 run or of more or less

chain-extended crystals with a thickness of 100 run and lateral dimensions of 10-12

Dm. Nevertheless, in both cases the lateral dimensions of these crystals are

extremely small compared to the lateral dimensions of solution-crystallized UHMW­

PE samples, which can be of the order of hundreds of nanometers up to several

micrometers [12,13].

Figure A.I Transmission electron micrograph of a nascent UHMW-PE reactor

powder embedded in epoxy

Figure A.2 shows the melting behaviour of the nascent reactor powder revealing a

peak melting temperature of 142.3 0c. The second run, after melting and re-

116 Appendix. ._------_ ......... _-_._------

crystallization, shows a significantly lower and broader melting endotherm with a

peak melting temperature of 133_0 "C_ Combining these DSC data with the electron

micro~copy re~u1t~, One is tempted to conclude that extended-chain crystals prevail

in these samples, in line with literature data [2,6,7,9J.

__ lSI (un

--_ ... ","' 2r"'1d run

------~~=-"-~~

~-

142.3

f)01-----,--~-·-·· .. T""·~~-1~OO-~-~-~-~-1·5-0-----i

- Temperature I'C)

Figure A.2 DSC them10grams of a nascent UHMW·PE reactor powder; first and

second run

However, an alternative way to explain the enhanced melting temperature of some

UHMW·PE reactor powders, i:s the possibility that during heating in the DSC fast

chain reorganizations take place during which the nascent crystals adapt constantly

and effectively to the applied temperature.

Assuming that the chain mobility upon heating is related to the entanglement

density and/or the stem arrangement within the crystals, a difference in rate of

reorganization upon annealing Or heating could be contemplated in favour of both

nascent reactor powders and solution-crystallized polyethylene in general, since

these materials p05:sess 11. low entanglement density as demonstrated by the high

drawability_

This concept is difficult to prove although remarkable segmental mobility has been

observed in the case of melting disentangled UHMW-PE samples [14,15}. An

Morphology of Nascent UHMW.PE 117

additional factor, promoting thermal instability in the case of nascent reactor

powders, is their small lateral crystal dimensions [16].

To prevent reorganization during heating in a DSC, the nascent reactor powders

were crosslinked uSing electron beam radiation. According to Bair et al. [17}, the

annealing of polyethylene single crystals can be hindered via crosslinking of the

amorphOllS phase through electron beam irradiation. They observed that poly"

ethylene single crystals, originally showing multiple peak melting, melted without

any reorganization after 260 kGy electron beam irradiation and without any signi­

ficant damage to the crystals.

The samples used in Our study were irradiated with a dose of 250 kGy. Afterwards,

the melting temperature of the reactor powder was 133.7 QC (see figure A3) which

is consiste.nt with the melting temperature expected for chain-folded lamellae

(compare with figure A.2, 2nd run).

-~ EB irrWiated (250 kGy)

----- nascent

50 100

142.3

~ (I ( I ( I I I I I J I I I

J I ( I ( I

I

\

150 ____ Temperature (OC)

Figure A.3 DSC thermograms of an UHMW-PE reactor powder; nascent and after

250 kGy electron beam irradiation

It might be argued that this drop in melting temperature after EB irradiation is not

due to hindered annealing but is caused by (partial) destruction of the crystals by

118 Appendix.

the electron beam. However, according to Keller {IS], the melting temper,tture

of mo:;t polyethylene samples is. virtua.lly unaffected up to 10,000 kGy EB

irradiation. In this study, a dose of 250 kGy was used, resulting in a decrease in

melting temperature of about 9 ~c. Furthermore, the crystallinity of the re,tctor

powder did not change significantly upon electron beam irradiation, which makes

radiation damage as a cause for the large drop in melting temperature rather

unlikely.

The observations described above, support our assumption that the high melting

temperature of certain nascent polyethylene reactor powders is not due to chain­

extended crystals but to fast annealing of chain-folded lamellae, raising the melting

temperature significantly during a DSC run. These chain-folded lamellae have a

thickness of about 10 nm and lateral dimensions of at most 100 nm.

A.4 Conclusions

Transmission electron micrographs of nascent polyethylene reactor powders

prepared under conditions yielding largely disentangled, highly drawable products

show, that these powders consist of relatively small (10 x 100 nm), irregularly

stacked crystals, which explains the absence of any SAXD intensity.

The nascent polyethylene powders possess a high melting point (142.3 ~C), related

to fast annealing of chain-folded crystals during a DSC run whereby the chain

reorganization is enhanced by the low entanglement density, the 'ideal' stem

arrangement and, in particularly, the extremely small crystal dimensions_

Electron heam irradiation can be used to crosslink the amorphou5 pha.se and

consequently limit the annealing process, resulting in a melting point of the reactor

powder of 133.7 ~c.

The results indicate that the commonly accepted relationship between a high mel­

ting point and the presence of -more or less- chain-extended crystals in polyethy­

lene reactor powders can not be generalized. Further experimental evidence suppor-

Morphology of Nascent UHMW-PE 119

ting the concept of a relation between chain reorganization rate, entangle-ment

density, stem arrangement and crystal dimensions, will be published in the future.

A.S References

L Smith, P., Chanzy, H.D. and Rotzinger, B.P., Polym, Comm. 1985,26,258

2. Smith, P., Chanzy, H.D. and Rotzinger, B.P.,l. Materials Sci. 1987, 22, 523

3. Ro~inger, RP., Chanzy, H.D. and Smith, P., Polymer 1989, 30, 1814

4. Smith, P., Lemstra, P J. and Booij, H.C,!. Polym. Sci., Polym. Phys. Ed. 1981,

19,877

5. Chamy, H.D" Revol, J.F., Marchessault, RH. and Lamande, A., Kolloid-2.

u. Z. Polymere 1973, 251,563

6. Chanzy, RD., Bonjour, E. and Marchessault, RH., Colloid & Polymer Sci.

1974,252,8

7. Kellerj A, and Willmouth, P.M., Die Makmmolekulare Chemie 1969, 121,42

8. Blais, P. and st. John Manley, R., J. Polym. Sci., Part A·] 1968, 6, 291

9. Chamy, H., Day, A. and Marchessault, R.H., Polymer 1967, 8, 567

10. Wunderlich, B., Hellmuth, E., Jaffe, M., Liberti, P. and Rankin, J., Kolloid.Z.

u. 2. Polymere 1965, 204, 125

11. Montezinos, D., Wells, B.G. and Burns, J.L, 1. Polym. Sci, Polym. Lett. Ed.

1985, 23, 421

12. Wunderlich, B., 'Macromolecular Physics, Volume 2: Crystal Nuc:leation,

Growth, Annealing', Academic Press, New York, 1976

13. Bastiaansen, C.W.M., Froehling, P., Pijpers, Al. and Lemstra, P.l., 'Integrati.

On of Fundamental Polymer Sc:ience and Technology', ed. by Kleintjens. L.A.

and Lemstra, P.J., Elsevier Applied Science Publishers Ltd., London, 1986,

p. 508

14. Bastiaansen, C.W.M., Meyer, H.E.H. and Lemstra, P.I., Polymer 1990, 31,

120 Appendix

1435

15. Barham, P.I., private communications

16. Wunderlich, R, 'Macromolecular Physics, Volume 3: Crystal Melting',

Academic Press, New York, 1980, chapter 8, p. 30

17. Hair, H.E., Salovey, R and Hll~eby, TW., Polymer 1967,8,9

18. Keller, A, 'Developments in Crystalline Polymers-l', ed. by Bassett, D.C.,

Applied Science Publishers Ltd., London, 1982, chapter 2, p. 81

Summary 121

Summary

Ultra-High Molecular Weight Polyethylene (UHMW-PE) fibres prepared via

solution-spinning and drawing, possess impressive shorHerm properties like a high

specific modulus and tensile strength and a large work to break. The use of these

fibres for applications in structural composites however, is limited due to some less

favourable properties like a relatively low melting temperature, creep under

prolonged static loading and relatively low shear moduli.

Both the positive and the less favourable properties of these fibres are governed by

their highly anisotropic nature_ Mter drawing, the fibres consist of oriented

crystalline regions consisting of more or less extended chains, alternated by

disordered domains containing physical entanglements, chain ends and interlamellar

tie-molecules. The strong covalent bonds in the fibre direction cause high stiffness

and strength levels but the weak intermolecular Van der Waals interactions

operative pe.:pendicular to the chain axes, cause the low creep resistance and off­

axis properties.

These properties are expected to be improved by increasing the degree of

intermolecular interactions between the chains. In the past, several methods have

been studied to introduce more interactions in the fibres. One of these routes

involves electron beam (EB) irradiation of the oriented stmcture with the aim of

introducing chemical crosslinks_ Unfortunately, EB irradiation of these highly

crystalline fibres not only causes crosslinking but also chain scissioning, which

results in decreases in tensile strength and creep resistance. The ratio

crosslinkingjscission might be improved by introducing unsaturated C-C bonds in

the system.

122 Summary

In this thesis the use of trans-l,4-polybutadiene (t-l,4-PB) in oriented structures is

explored, with respect to ED irradiation. T-l,4-PB is a linear polymer containing

one unsaturated C-C bond per 4 C-C bonds, with a high tendency to crosslink. Fur­

thermore, it has a conformationally disordered crystal phase between 70 "c and 140

OCt in which the molecular segments are mobile. 111is mobility is expected to en­

hance crosslinking upon irradiation ba.~ed on similar experiments with paraffines.

Oriented structures of pure t-1,4-PB can be obtained via solid-state coextrusion, a

rather academic technique, or by tensile drawing of irradiated films in which the

network facilitates stress transfer between the separate chains. Via both techniques,

the optimum properties are limited, a maximum Young's modulus and tenacity of

about 20 OPa and 0.4 OPa respectively can be attained. Post-irradiation of the

oriented t-l,4-PB structures does not improve these properties. Since the ultimate

properties of t-1,4-PB tapes are limited by the relatively low molecular weights used

and the unfavourable non-planar chain conformation, it seems interesting to blend

t-l,4-PB with UHMW·PE to combine the impressive fibre properties of UHMW-PE

with the EB susceptibility of t-l,4-PB.

Blending UHMW-PE and t-l,4-PB, with the aim of preparing blended fibres in the

future, via the gel-spinning technique, should take place in solution in order to

obtain largely disentangled, ultra-drawable precursors. In this thesis, blending of

both components was performed in 1.5 wt.% solution:s in xylene. It appears that

crystallization-induced pha:se separation occur:s, re:sulting in hlended films consisting

of both UHMW-PE and t·l,4-PB crystals. Probably some miscibility of both

components occurs in the disordered domains due to simultaneous solidification of

amorphous PE and t-l,4-PB upon evaporation of xylene.

Via uniaxial tensile drawing of the blends at 120°C, draw ratios up to 170 can be

reached, depending on the t-l,4-PB content of the blends. The morphology of the

blends changes drastically upon drawing. The PE domains in the blend show the

same behaviour as encountered upon drawing of pure lJHMW·PE. The t-1,4-PB

Summary 123

crystals show the same orientation behaviour as the PE crystals in the initial stages

of drawing. Further drawing results in an elongation of the PB crystals by the

deforming PE matrix, leading to elongated t"I,4-PB domains with a thickness of

only a few nanometres. The fraction monoclinic t-1 ,4-PB in these domains decreases

with increasing draw ratio in favour of hexagonal and/or amorphous.

As expected, the final tensile properties of the blends (modulus and tensile

strength) decrease with increasing t-IA-PB content

EB irradiation of the drawn blends results in the fonnation of a network. Both the

gel fraction and the network density increase with increasing irradiation dose and

irradiation temperature. The presence of t" 1,4"PB in the tapes lowers the gel point

dose of the system and seems to reduce the amount of chain scission as appears

from a smaller decrease in tensile strength upon irradiation of blends compared

with pure polyethylene. The Young's modulus of the blends is hardly inftuen(::ed by

EB irradiation.

The network formed upon irradiation of the drawn blends, can sustain constrained

heating to 200 °C, in contrast to irradiated pure PE tapes which melt and fail under

these conditions. It is however, not capable of improving the creep resistance and

off-axis properties of the drawn blends.

Based on all observation, a model can be derived for the network structure jn the

drawn blends. This structure consists of elongated t-t,4-PB domains finely dispersed

in a matrix consisting of PE fibrils. These PE fibrils contain disordered domains

(defect zones) jn which some t-l,4-PB chains are trapped and mixed on a molecular

scale. The crosslinks situated in these disordered domains provide for the continuity

of the network but are not sufficient to improve the creep resistance. For that

purpose probably crosslinking within the crystalline regions is necessary.

Since the morphology of the oriented system appears to determine the crosslinking

efficiency, it is difficult to transfer the results obtained in this thesis with tapes

prepared from films, directly to fibres prepared via solution·spinning.

Samenvatting 125

Samenvatting

Ultra-Hoog Molecuulgewicht Polyetheen (UHMW-PE) vezels, gemaakt via gel­

spinnen en verstrekken, vertonen indrukwekkende eigenschappen als een hoge

modulus en tteksterkte en een hoge energie-absorptie. Ten gevolge van eel1 aantal

minder gunstige eigenscbappen, een lage temperatuurresistentie, kruip onder

constante belasting en lage afschuifmoduli, worden de toepassingsmogeIijkheden in

structurele composieten bepetkt.

Zowel de positieve all! de minder gunstige eigenschappen van de vezels zijn

inherent aan de extreem anisotrope sttuctuur van verstrekt polyetheen. Na verstrek­

ken bestaat de stIUctuur uit georienteerde kriStallijnot; gebieden afgewisseld met

ongeordende (amorfe ) zones die fysische entanglements, keteneinden en

interlamellaire verbindingsmoleculell bevatten. De sterke covalente bindingen in de

vezelrichting wrgen voor de hoge modulus en treksterkte terwijl de zwakke Van der

Waals interacties, die werkzaam zijn tussen de ketens, de lage kruipweerstand en

transversale eigenschappen veroonaken. Deze laatste eigenschappen zouden

mogelijkerwijs verbeterd kunnen worden door het introduceren van additionele

intermoleculaire interacties. In het verleden zijn hiervoor verschillende methodes

bestudeerd. Ben daarvan omvat elektronenbestraling van de vezels met als doel het

introduceren van cheInische knooppunten (crosslinks). Helaas gaat elektronen­

bestraling met alleen gepaard met vernetting (crosslinking) maar ook met

ketenbreuk, hetgeen resulteert in een afname van de treksterkte en de kruipweer­

stand. De vcrhouding tussen crosslinking en ketenbreuk zou misschien verbeterd

kunnen worden door onverzadigde bindingen in het systeem te introduceren.

126 Samenvatting

In dit proefschrift worden de mogelijkbeden met betrekk..ing tot elektronenbestraling

onderzocht van georienteerde structuren gebaseerd op trans-l,4-polybutadieen.

T-l,4-PB heeft een zekere gelijkenis met polyetheen, beide zijn Iineaire semi­

kristallijne polymeren met een smeltpunt van ongeveer 140°C. T-l,4-PB beeft

echter een onverzadigde C-C binding per vier C-C bindingcn en bezit daardoor een

grote neiging tot crosslinken. Bovendien vertoont het een 'condis' (conformationally

disordered) kristallijne fase tussen 70 en 140 °e, waarbij de molecuulsegmenten

zeer beweeglijk zijn. Deze hoge beweegJijkbeid zou het vernetten positief kUlll1en

belnvloeden zoals het gevaJ is bij het bestraten van parafines in een vergelijkbare

'disordered' fase.

Puur t-1,4-PB kan georienteerd worden door middel van vaste-stof-coextrusie, of via

uniaxiaal verstrekken van bestraalde films waarbij de gevormde crosslinks

spanningsoverdracht tus~en de verschillende ketens mogeJijk maken. De optimale

eigenschappen die mct beide technieken verkregen kunnen worden zijn beperkt:

een maximale modulus en treksterkte van respectievelijk 20 GPa en 0.4 GPa. Na­

bestralen van de georienteerde t-l,4-PB monsters levert geen verbetering van deze

eigenschappen op. Gezien de relatief lage maximaal haalbare eigenschappen van

georieenteerde t-l,4-PB structuren, inherent aan de niet planaire ketenconformatic

en de lage beschikbare molmassa's, lijkt het interessant om t-l,4-PB en OHMW.PE

te mengen en zodoende de vezeleigenschappen van PE te corobinercn met de

gevoeligheid van t-l,4-PB voorelektronenbestraHng.

Het mengen van PE en t·l,4.PB voor het maken van vezels moet plaatsvinden in

oplossing am hoog-verstrekbare systemen te verkrijgen. In het onderzoek beschre­

vcn in dit proefschrift werd het mengen uitgevoerd in 1,5 gew.% oplossingen in

xyleen. Het blijkt dat door afzonderlijke kristallisatie van beide componenten,

fasenscheiding geYntroduceerd wordt, hetgeen resulteert in films waarin zowel PE

kri~tal1cn ats t-l,4-PB kristallen aangetoond kunnen worden. Waarschijnlijk zijn de

beide polymeren in gering mate gemengd in de ongeordende (amorfe) gebieden in

de films, ten gevolge van geIijktijdige solidificatie van amorf PE en t-1,4.PB.

Samenvatting 127

Via uniaxiaal verstrekken van de blends bij 120°C lrunnen verstrekgra.den tot 170

bereikt worden, afbankelijk van het t-l,4-PB gehalte. De morfologie van de blends

verandert drastisch tijdens het verstrekken. De PE gebieden vertonen hetzelfde

gedrag als bij het verstrekken van puur PE. De t-1,4-PB kristallcn vertonen injtieeJ

hetzelfde orientatiegedrag als de PE kristallen maar worden bij verder verstrekken

vervormd door de deformerende PE-matrix, hetgeen resulteert in lange, dunne t­

l,4-PB gebieden met een dikte van enkele nanometers. De fractie monoclien PB

in deze gebieden neemt af met toenemende verstrekgraad ten gunste van hexago­

naaJ en/of amorf.

Zoals verwacht nemen de mechanische eigenschappen van de blends (modulus en

treksterkte) af met toenemend t-1,4-PB gehaJte.

Elektronenbestraling van de verstrekte blends resulteert in netwerkvorming. Zowel

de gelfractie als de netwerkdichtheid nemen toe met hogere bestralingsdoses en

bestralingstemperaturen. De aanwezigheid van t-I,4·PB in de monsters verIaagt de

gelpuntsdosis en het optreden van ketenbl"euk; het laatste blijkt uit een geringere

afname van de treksterkte van blends door bestralen dan van puur polyetheen_ De

modulus van de blends wordt nauwelijks be"invloed door de bestraling.

Het gevormde netwerk in de blends blijft intact als de monsters (op constante

lengte gehouden) opgewarmd worden tot 200 QC, in tegenstelling tot vergelijkbare

pure PE monsters die smelten en breken onder deze condities_ Het netwerk is

echter Diet in staat de kruipweerstand en transversale eigenschappen te verbeteren.

De waarncIllingen kunnen verklaard worden uitgaande van een modelvoorsteHing

van de netwerkstructuur in de verstrekte blends. Deze structuur bestaat uit een

matrix van PE fibriUen waartussen op fijne schaal, dunne uitgerekte PB domeinen

gedispergeerd zijn. Verder 2;ijn in de ongeordende zOnes (defectzones) in de PE

fibrillen in geringe mate t-l,4-PB keten..~ ingevangen, die op moleculaire schaal

gemengd zijn. De crossHnks die in deze ongeordende domeinen gevormd worden

zijn verantwoordelijk vOor de continuiteit van het netwerk maar met voldocnde vOOr

128 Samenvatting -------_., .... ------

een verbetering van de kruipweerstand. Daarvoor is waarschijnlijk vernetting van

de kristallijne fase noodzakelijk.

Gezien het beJang van de morlologie van de verstrekte bJends voor de effectiviteit

van de te vormen netwerken, is het niet zonder meer mogelijk de resultaten be­

~chreven in dit proefschrift, te extrapo!eren naar vezels geproduceerd via het gel­

spin proces.

Nawoord 129

Nawoord

Op deze plaats wil ik iedereen, die op enigedei wijze beeft bijgedragen aan het

ondeIZOek en het tot stand komen van dit proefschrift, hartelijk bedanken. Een

aantal personen wil ik daarbij met name noemen;

• Mijn co-promotor Cees Bastiaansen (DSM Research). die het onderzoek

gelnitieerd en enthousiast begeleid heeft.

• Prof. Moller van de Universiteit Twente (en zijn voormalige medewerkers

Horst Deckmann en Martin Kunz) voor vele discussies en adviezen op het

gebied van NMR en TEM.

• De heren de Haan en van de Yen voor NMR metingen en hulp bij de inter­

pretatie daarvan.

• Marinus Hom van het IRI in Delft vOOr het uitvoeren van de bestralingen

en Erik-Jan Langkaro.p voor de analyse van de bestraalde monsters.

Alle collega's en ex-roll ega's van de vakgroep TPK vOor de prettige

samenwerkin& met name Mark Saveisberg voor de synthese van trans-l,4-

polybutadieen en Leon Govaert voor zijn hulp bij de mcchanische

experimenten.

Curriculum Vitae 131

Curriculum Vitae

Naam Yvonne Tervoort-Engelen .

~ 71-Geboren 13 december 1964 te Nijmegen

Opleiding OngedeeJd VWO, 1977-1983, Canisius-College Mater Dei

te Nijmegen

• le fase Scheikundige Technologie, 1983-1987, Technische

Universiteit Eindhoven

• Onden::oekersopleiding Scheikundige Technologie, 1987-

1989, Instituut Vervolg Opleidingen, TUE

Assistent in Opleiding Scheikundige Technologie, 1989·

1991, Vakgrocp Polymeerchemie en Kunststoftechnologic,

TUE

Werkzaam Onderzoekspeciali5t elektronenmicroscopie, vanaf 1 september

1991, Vakgroep Polymcerchemie en KUl1ststoftechnologie, TUE

Stellingen

L Het verbeteren van de kmipweerstand van UHMW-PE vezels is gelimiteerd

door het feit dat de tot nu toe toegepaste methoden de kristallijne fase van

de vezels intact laten.

Chen, YL. and RJ.nby, B., Polym. Adv. Techn. 1990, L 103; Dit proef~chrift,

hoofdstuk 1 en hoofdstuk 6

2. Ottam en Porter hebben bij de interpretatie van hun calorimetrisch en

RAMAN/LAM onderzoek van UHMW·PE reaktorpoedcrs, onvoldoende

rekening gehouden met de mogeJijkbeid van reer snelle kristaireorganisaties,

die kunnen optreden in k1eine instabiele kristallen die gevormd worden

tijdens poiymerisatie bij lage temperaturen.

Dttan;' S. and Porter, RS., J. Polym. Sci. Polym. Phys. Ed. 1991, .22 1179;

Ottani, S. and Porter, RS., 1. Polym. Sci. Polym. Phys. Ed. 1991, Z2 1189; Dit

proefschrift, appenda

3. De door Yamaura et al. gerapporteerde waarden voor de Young's modulus

en de treksterkte van polyvinylaIcohol vezels zijn onwaarsc;hijnlijk hoog.

Yamaura, K, Tanigami, T., Hayashi, N., Kosuda, K, Okuda, S., Takemwa, Y.,

Itoh, M. and Matsuzawa, S" J. Appl. Polym. Sci. 1990, 1f), 905

4. De relatie tussen de compliantie in de stationaire toestand, de nul-viscositeit

en de eerste normaa]spanningscoefficient wordt door Wolf ten onrechte

toegepast in het niet-lineaire visco-elastische gebicd.

Wolf, B.A., Macromolecules 1984, 12 615; Bird, R.B., Annstrong, R.C and

Ha.~sager, 0., 'Dynamics of Polymeric Liquids', volume 1: Fluid Mechanics,

John lViley & Sons Inc., New York, 1987, p. 130

5. De kristallisatietheorie voor polymeren is nog lang niet uitgekristalliseerd.

Dit pmefschrift hoofdstuk 3 en appendix

6- De door Stehling et al. waargenomen omslag van een 'core-shell' naar een

'inter,penetrating' structuur van de disperse fase In ternaire

PP !HDPE!EPDM blends met toenernende HDPE!EPDM verhouding, kan

ook door verschillen in verwerkingscondities optreden.

Stehling, F.C, Hufh T., Speed, C.S. and WISSler, G., 1. Appl. Polym. Sci. 1981, 2Q. 2693; Tervoort-Engelen, Y.M.T. and van Gisbergen, J.G.M., Polym. Comm. 1991, n.. 261

7. De relatief lage dichtheid van poly-B-hydroxybutyraat (PHB) granules in de

Alcaligenes eutrophus bacterie, wordt in tegenste1lling tot de bewering van

Mas et aL, niet veroorzaakt door de aanwezigheid van 40 % water maar

door het feit dat het aanwezige PRE amorf is.

Mas, 1-, Pedros-Alio, C. and Guen'ero, R, 1. Bacteriology 1985, l.Q1. 749;

Barham, P.J., Keller, A., Glun, E.L. and Holmes, P A, 1. Mater. Sci. 1984, l!l 2781; Barnard, G.N. and Sanders, 1-KM., 1. BioI. Chem. 1989, 2Q1, 3286

8. De director-orienta tie van de moleculen in vloeibaar kri~;tallijne

diacrylaatfilms waarin de moleculaire ordening is gefixeerd door UY­polymerisatie, kan afgeleid worden ult de lagenstructuur die met SEM

waargenomen wordt in breukvlakken van deze films.

Heynderic/o;, i., Bro(~r, D']. and TelVoort-Engelen, Y.M. T., accepted for

publication in 1. Mater. Sci.

9. De eenstaps-synthese route van 2-(3H)-benzofuranonen m.b.v. trifluor­

azijnzuur, zoaJs recentelijk beschreven door Chaturvedi et al., is reeds

gangbaar in de literatuur en de:rhalve het vermelden in een communication

niet waard.

Chaturvedi, R and MulchandQ.m~ N.N., Synthetic Comm. 1990, 2Jl 3317;

March,l., 'Advanced Organic Chernistry~ 2nd. Ed., McGraw-Hill, Auckland

10. Naast een 'botcrberg' en cen 'melkplas' bestaat er ook een 'publicatiezee'.

Eindhoven, 17 december 1991 Yvonne Tervoort-EngeIen