Characterisation Of Polymers By Thermal Analysis - W Groenewoud (Elsevier, 2001) Ww

CHARACTERISATION OF POLYMERS BY

THERMAL ANALYSIS

This Page Intentionally Left Blank

CHARACTERISATION OF POLYMERS BY

THERMAL ANALYSIS

W.M. GROENEWOUD

Eerste Hervendreef 32, 5232 JK 's Hertogenbosch The Netherlands

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First edition 2001 Second impression 2003

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Trans fe r red to digital pr in t ing 2005

To Vera for 36 years of love, support and continuous inspiration.


PREFACE

The development of the Linear Variable Displacement Transducer (LVDT) was a first order technological break-through after centuries in optical length-difference measurments. The first LVDT's became commercially available in Holland in 1959. Our research team (I was a junior member) bought one LVDT for the development of a length dilatometer to measure the change in length of a polymer sample during a heating or cooling procedure. The LVDT signal and (sample temperature) thermocouple signal were recorded on an XY-recorder. Indeed, we were very proud of our first 'automated' measuring system. We did not yet call our system a Thermal Mechanical Analyser (TMA) nor described our activities as 'Thermal Analysis'. Nowadays, computer controlled dynamic and static TMA systems are supplied by several manufacturers and perform completely automated the measuring and data handling procedures required.

This story illustrates the huge technological development during the last forty years. Thermal Analysis (TA) has become an indispensable family of analytical techniques in polymer research. This increased importance of these techniques can be seen as the result of three more or less parallel developments: - a tempestuous development of TA measuring techniques in

combination with a high degree of automation, - the strongly increased understanding of the underlying

theory, published by authors like Wunderlich, Hohne, Richardson and Mathot [1,2] and

- the increasing knowledge of the relation between the polymers' chemical structure and their physical properties.

These developments still continue and a lot of work has yet to be done in the second and especially the third area. Increasing knowledge of the dependence of physical properties on chemical structure form the added value of accurate thermoanalytical measurements and this knowledge is very important for the development of new polymeric systems.

The table below lists the various TA techniques following the notation of the ICTA (International Committee for Thermal Analysis) nomenclature committee. The three "classic" TA techniques are DSC, TGA and TMA of which DSC is still the "workhorse". TA is also covering, however, a substantial number of other techniques and applications and several of these techniques are described in this book. This book is not a comprehensive textbook about TA but more a survey of the author's work during many years, at the Koninklijke Shell Laboratorium in Amsterdam. It describes in six chapters the use of the various TA techniques (printed in bold in the table) for specific problems, illustrating the versatility of TA. A technical description is only given for equipment of own design.

Thermal Analysis techniques

Differential Scanning Calorimetry (DSC) - high pressure DSC - photo-DSC - modulated DSC

Thermogravimetric Analysis (TGA)

Thermodilatometry - length dilatometry (TMA) - volume dilatometry

Dynamic Mechanical Analysis (DMA) - s t a n d a r d (low frequency) DMA -ultrasonic (high frequency) analysis

Thermo-electrometry - dielectric analysis (DETA) - volume resistivity analysis - thermally stimulated discharge analysis

Simultaneous Techniques - thermally stimulated discharge analysis/thermomechanical

analysis (TSD/TMA) - thermogravimetric analysis/fourier transform infra red/mass

spectroscopy (TGA/FTIR/MS) - thermogravimetric analysis/differential thermal analysis/

mass spectroscopy (TGA/DTA/MS) - thermomechanical analysis/dielectric analysis (TMA/DETA)

Thermoluminosence Thermomagnet ome try Thermo-optometry Thermosonimetry

Over the years, quantitative structure/property relationships have been developed by various workers in the polymer research field. Well known are for example the important contributions made by D.W. van Krevelen in 'Properties of polymers' [3] and by J. Bicerano in 'Prediction of Polymer Properties' [4]. An endeavour is made in chapter seven (and eight) to improve some of such correlations by using consistently measured, reproducible TA data. Chapter nine shows the contribution of TA to the characterisation effort necessary for the technical and commercial development of a new polymer system. Chapter ten finally, provides information about different polymers obtained during special case studies. This book illustrates in this way, applications of a wide variety of thermal analysis techniques. The author hopes that this monograph will be useful especially to those who are going to work in the thermal analysis area in the context of polymer research.

Wire Groenewoud

ACKNOWLEDGMENTS

The results described in this book could only be obtained by the expertise and the cooperation of many members of the different skillgroups at the Koninklijke Shell Laboratorium in Amsterdam (KSLA). The still unique possibilities of this laboratory are mentioned with pleasure. Without pretending to be complete, I have to mention a number of colleagues: For stimulating discussions and valuable insight provided by Roel Jongepier, Bram Ghijsels, Toni Cervenka, John Wintraecken and Piet Kooijmans. An important part of the experimental work was performed by: Arie van der Zwan, Nico Groesbeek, Ton Jakobs, Wouter de Jong, Bob Oudhaarlem, Leo Sman and Karin Orzessek. Bob van Wingerden read and discussed with me many of the internal reports which formed the basic data source of this book resulting in many, always improving, suggestions. Regretting any unintentional omissions I finally thank the management of KSLA for the permission to publish results of our polymer research.


CONTENTS

Preface

Acknowledgments

Chapter 1. Differential Scanning Calorimetry I. 1 Introduction

I.I.I DSC calibration and stability 1.1.2 The Tg-value determination 1.1.2 Melting/recrystallisation determinations

10 Ii 14

1.2 Tg-values of car-tyre rubber systems 1.2.1 Introduction 1.2.2 The Tg-value of BR and SSBR rubbers 1.2.3 The Tg-value after blending and oil-extension

of BR and SBR rubbers

17 17

19

1.3 Recrystallisation and fusion of polypropylene 1.3.1 Introduction 1.3.2 Additives acting as nucleating agents for PP 1.3.3 Annealing experiments with i-PP

26 26 28

1.4 Side-chain crystallisation in poly(1-olefine) s 1.4.1 Introduction 1.4.2 Crystallisation in poly(l-olefin)s

36 36

1.5 Chemical reactions monitored by DSC I. 5.1 Introduction 40 1.5.2 The determination of the cure conditions of

a powder coating system 43 1..5.3 Reactions of model compounds studied by DSC 43

1.6 Determination of the heat of vaporisation by DSC 1.6.1 Introduction 52 1.6.2 DSC modification for the AHvap.25 determination 52 1.6.3 Results of AHvap.25 determinations by DSC 54

Chapter 2. Thermogravimetrical Analysis 2.1 Introduction 61

2.2 01igomers content and thermal stability of polypropylene 2.2.1 The non-isothermal thermal stability

determination 63 2.2.2 The isothermal thermal stability determination 65

2.3 The TG analysis of a PP catalyst system 2.3.1 A 'plastic wrapped' TGA 70 2.3.2 TG analysis of a MgCl2-supported, TiCI4/DIBP

catalyst 72

Chapter 3. Thermodilatometry 3.1 Length dilatometry (TMA)

3. i. 1 Introduction 3.1.2 The l.e.c, determination of filled polymers 3.1.3 Shrinkage of PK terpolymer and nylon 6.6 due

to moisture loss

3.2 Volume dilatometry 3.2.1 Introduction 3.2.2 The volume dilatometer 3.2.3 The measuring procedure 3.2.4 Isothermal crystallisation of IR rubbers

77 77

81

85 85 89 91

Chapter 4. Dynamic Mechanical Analysis 4.1 The standard DMA technique

4.1.1 Introduction 4.1.2 DMA analysis of PP/C2C3 rubber blends 4.1.3 Tg-value determination of aged, rigid PU

foams by DMA

4.2 Mechanical measurements at ultra-sonic frequencies 4.2.1 Introduction 4.2.2 The ultra-sonic measuring equipment 4.2.3 Results of ultra-sonic measurements on

car- tyre rubbers

94 99

105

109 112

114

Chapter 5. Thermo-electrometry 5.1 The DC and AC properties of polymers

5.1.i Introduction 5.1.2 DC properties of polymers 5. i. 3 AC properties of polymers

123 124 128

5.1.4 The AC and DC measuring system 132 5.1.5 AC and DC properties of a cured resin system 134 5. i. 6 Time/temperature superposition of dielectric

results 140 5.1.7 The dielectric constant of rigid PU foam 145

5.2 Effect of moisture on the electrical properties of polymers 5.2.1 Introduction 5.2.2 Influence of moisture on the dielectric

properties of resin castings and laminates 5.2.3 Effect of seawater and cargo on the electrical

properties of a tankcoating system 5.2.4 The determination of the Ki-value of PVC cable

insulation

151

153

158

163

5.3 Conduction improvement of epoxy resins by carbon black addition 5.3.1 Electrostatic safety criteria 171 5.3.2 DC properties of experimental epoxy resin/

carbon black systems 172 5.3.3 DC properties of anti-static epoxy GFR pipes 177

5.4 Thermally Stimulated Discharge analysis 5.4.1 The TSD technique 5.4.2 Bucci's TSD theory 5.4.3 Results of TSD experiments

181 181 184

Chapter 6. Coupled thermal analysis techniques 6.1 Introduction

6.2 Simultaneous TSD/TMA measurements 6.2.1 The TSD/TMA system 6.2.2 TSD/TMA results

6.3 The TGA - coupled - FTIR/MS technique 6.3.1 Introduction 6.3.2 The TGA/FTIR and TGA/MS coupling 6.3.3 The heated capillaries tip temperatures 6.3.4 Single component calibration 6.3.5 Investigation of the thermal decomposition

of Cobaltphthalocyanine by TGA - coupled - FTIR/MS

6.3.6 Investigation of the released vapours during the cure of epoxy resin system by TGA - coupled - FTIR/MS

188

191 192

195 196 200 201

209

222

Chapter 7. Chemical structure/physical properties correlations

7.1 Introduction 230

7.2 The Tg-value estimation 7.2.1 Introduction 232 7.2.2 The 'modified cohesion energy' method 233 7.2.3 The Tg-value of crosslinked polymeric systems 245

7.3 The Tm-value estimation 7.3.1 Introduction 7.3.2 The reduced Tg/Tm correlations

253 254

7.4 The Hf-value estimation 264

7.5 The thermal stability estimation 7.5.1 Introduction 7.5.2 The semi-static Td, o-value determination 7.5.3 Thermal stability estimation based on

Td, o-values

268 269

269

7.6 The moisture sensitivity estimation 274

7.7 Estimation of the key-properties of a new polymer 277

Chapter 8. Tg-values of polymers with double bonds in the main chain and Tg-values of non-polar polymers with side chains


8.2 Experimental BR systems 8.2.1 BR with a high 1,4 trans content 8.2.2 BR with a high syndiotactic 1,2 BR content

282 286

8.3 Experimental IR systems 288

8.4 A Tg/structure correlation for non-polar polymer systems with side-chains 293

Chapter 9. Characterisation of polyketone polymer systems by Thermal Analysis Techniques


9.2 Investigation of the crystalline phase of PK co- and terpolymers 9.2.1 PK copolymer and PK terpolymer 297 9.2.2 The Tm(o)- and Hf(max.)-values of PK copolymer 299 9.2.3 Alpha- and beta-crystallinity in PK copolymer 302 9.2.4 Alpha- and beta-crystallinity in PK copolymer

after a common processing procedure 308 9.2.5 Alpha- and beta-crystallinity in PK terpolymers 310

9.3 Investigation of the amorphous phase of PK terpolymer by DMA/DSC 9.3.1 Amorphous phase transition effects 9.3.2 Ageing and moisture absorption effects 9.3.3 Determination of the Tg-value of PK terpolymer

by DSC

312 312

318

9.4 TMA measurements on PK terpolymer systems 9.4.1 The linear thermal expansion coefficient of

long glassfibre reinforced PK terpolymers 9.4.2 The repeatability of the l.e.c, determination

on PK terpolymer systems

322

325

9.5 Determination of electrical properties of PK terpolymers 9.5.1 The influence of moisture on the dielectric

properties 9.5.2 The frequency dependency of the dielectric

properties 9.5.3 The specific volume resistivity determination

of PK terpolymer

327

331

334

9.6 Survey of PK terpolymer thermal analytical characterisation results 337

9

Chapter i0. Thermo-analytical case studies

i0.i Introduction

10.2 The effect of the presence of a solvent during the cure of a thermoharding system

10.3 The thermal transitions of a liquid crystalline polymer

10.4 The optimal crystallisation temperature of diphenylolmethane

10.5 The dynamic stiffness of ultra-high molecular weight polypropylene in its melt

10.6 The effect of an anti-static additive on the electrical resistivity of a polystyrene foam

10.7 The dielectric constant of polyethylene foil

10.8 The volume resistivity of epoxy based moulding powder systems during immersion in hot water

10.9 The determination of the composition of a cartyre rubber

I0.I0 The thermal stability of ASB

I0.ii The thermo-analytical characterisation of a maize based, 'green' polymer

339

339

342

345

350

354

356

359

364

366

371

Index 377

DIFFERENTIAL S CANNING CALORIMETRY

CHAPTER 1

I0

CHAPTER i: DIFFERENTIAL SCANNING CALORIMETRY

i. 1 Introduction

1 , 1 , 1 The D$C Differential scanning calorimetry is, according to the ICTA I nomenclature committee, a technique in which the heat flux (power) to the sample is monitored against time or temperature while the temperature of the sample, in a specified atmosphere, is programmed. In practice, the difference in heat flux to a pan containing the sample and an empty pan is monitored. The instrument used is a differential scanning calorimeter or DSC. The DSC is commercially available as a power-compensating DSC or as a heat-flux DSC.

The power-compensating DSC has two nearly identical (in terms of heat losses) measuring cells, one for the sample and one reference holder. Both cells are heated with separate heaters, their temperatures are measured with separate sensors. The temperature of both cells can be linearly varied as a function of time being controlled by an average-temperature control loop. A second-differential-control loop adjusts the power input as soon as a temperature difference starts to occur due to some exothermic or endothermic process in the sample. The differential power signal is recorded as a function of the actual sample temperature.

One single heater is used in the heat-flux DSC to increase the temperature of both the sample cell and the reference cell. Small temperature differences occurring due to exothermic/ endothermic effects in the sample are recorded as a function of the programmed temperature. Both systems are extensively described in the literature, more recently by Wunderlich [i].

The DSC is used (after proper calibration, see 1.1.2) in polymer research for mainly three different types of experiments. a) glass-rubber transition temperature (Tg-value)

determinations, see 1.1.3, b) melting/recrystallisation temperature and heat (Tm/Tc-

value and Hf/Hc-value) determinations, see 1.1.4, c) measurements on reacting systems (cure measurements). An example of monitoring chemical reactions by DSC is given in 1.5. Besides, the use of the DSC for a specific non-standard application is described in 1.6.

1.1.2 DSC calibration and stability The DSC measurements reported in this book are performed with the power-compensating DSC-2 and DSC-7 systems from Perkin Elmer. The block surrounding the DSC sample holders is kept at -150~ • I~ with the aid of a controlled liquid nitrogen supply, both cells are purged with helium (60 ml/minute). The standard temperature calibration is performed at a heating rate of 20~ using the melting effects of cyclohexane

~ICTA, International Committee for Thermal Analysis

Ii

(-87.06~ and 6.3~ indium (156.60~ and tin (231.88~ The computer controlled two point calibration procedure is performed using the cyclohexane -87.06~ value and the indium 156.60~ value. The heat of fusion of indium (Hf-value = 28.45 J/g) is introduced to perform the energy calibration. A tin sample fusion measurement is, subsequently, performed to check the possible deviation in the upper part of the temperature region. This deviation proved always to be less than 0.5~

If the temperature region of interest is ranging from about 100~ up to 350~ the two-point calibration procedure is performed using indium and tin. The melting effect of lead (327.4~ is used in that case as the high temperature check. The temperature and energy calibrations of the DSC-2 and DSC-7 are surprisingly stable as shown by a series of fusion measurements on the same indium sample placed in the F~_~_~DSC- 7 apparatus, see Table I.I. The average indium T(onset) value proved to be 156.6~ • 0.1~ whereas the average Hf-value proved to be 28.5 J/g • 0.3 J/g measured over a period of about three month while the system was in use five days a week.

Table i.I Results of a temperature calibration stability test of a Perkin Elmer DSC-7

time, days

0 ,

9

17

3O

91

,, ~,J ,j,, ,,~

T (onset), oc

156.59

156.54

156.75

156.47

156.47 156.53

deviation, Hf-value, ~ J/g

, . . . . .

-0.01 28.10 . . . . . . . . . .

-0.06 28.63 . . . . . .

+0.15 28.86 . . . . . . . . . . .

-0.13 28.40 . . . .

-0.13 28.40 -0.07 28.40

, , :,,,, , t,, ~ , , I',I, .... ,', . . . .

a. DSC cell base at -150~ b. Helium purge gas, 60 ml/minute, c. Indium sample 5.81 mg.

deviation

-0.35

+0.18 , ,,

+0.41 ......

-0.O5 , , I

-0.05 -O.O5

........ ~ , , ,~ ~ , .......

d. Indium T(onset) and Hf-values measured at 20~ second heating scan values after a first heating/ cooling scan between 120oc and 160~

1.1.3 Tu,valuedetermination The DSC is widely used to measure the glass-rubber transition temperature (Tg-value), which is an important parameter for polymer characterisation. The Tg-value represents the temperature region at which the (amorphous phase) of a polymer is transformed from a brittle, glassy material into a tough rubberlike liquid. This effect is accompanied by a 'step-wise' increase of the DSC heat flow/temperature or specific heat/ temperature curve. Enthalpy relaxation effects can hamper the

12

HEAT FLUX

OLAS'~'

u I ! 1r ILl z I,,- o K W ~ - =|

-~0 50 g o

Figure 1.1 The Tg(onset)-value

(3ZO)

( ) = COOLING RATE, ~ THROUGH Tg AFTER PREHEATING AT 150~

[ ] = AGEING TIME, DAYS AGEING AT ROOM TEMPERATURE FOLLOWING QUICK COOLING (320~ FROM 150~

I 51

(40) / Sl

(10) ~ . ~ s t

I 5 2

(2.5)

( 0 . 6 2 )

[03

1:23

10 5 0 g o 10 5 0 gO T E M P [ R A T U R E , ~

Figure 1.2 Effects of cooling rate and ageing time (heating rate 20~

13

DSC Tg-value determination. A standardised procedure is therefore necessary, to arrive at reproducible results.

In general two types of DSC thermograms can be obtained for the glass transition of a rigid polymer. Figure I.I shows these two types for a linear epoxy resin. The upper thermogram was obtained by scanning a sample of this resin at a rate of 20~ without pretreatment.

The lower thermogram was obtained on the same sample, which was preheated at 150oc for one minute, then quickly cooled (320~ to room temperature before scanning it under the same conditions. In the lower curve the glass transition is visible as the expected 'step-wise' heat flow shift. In the upper curve, however, a strong endothermic effect is superimposed on the heat flow shift. The temperature at the intersection of the extrapolated heat flow curve at the low temperature end and the tangent of the ascending curve at the inflection point is defined as Tg-value often indicated as DSC Tg(onset)-value. It is evident from Figure i.I the this Tg definition leads to different results for both thermograms. This is caused by the different thermal histories of the samples, which results in a difference in the extent of the so-called enthalpy relaxation effect [2].

Figure 1.2 illustrates, using the same sample, how the rate of cooling through Tg and storage at room temperature bring into evidence the presence of the enthalphy relaxation effect as a superposition on the heat flow curve shift. Figure 1.2 also shows the extent of the Tg(onset)-value differences due to the presence of these endothermic peaks. It will be clear that a standardised Tg-value determination procedure is necessary to obtain reproducible results-

- the sample (I0 to 15 mg.) is placed in the DSC sample holder,

- the sample is heated at a rate of 20~ through the possible present enthalphy relaxation maximum,

- the sample temperature is decreased, subsequently, at maximum cooling speed to a temperature of at least 50~ below the measured Tg effect,

- the sample is heated a second time through its Tg region at a rate of 20~ and this second scan result is used to calculate the Tg(onset)-value,

- the sample weight is checked to see if any weight loss occurred due to the thermal treatment of the sample (for instance due to loss of moisture).

The Tg-values reported in this book are measured according to this procedure. A series of Tg(onset)-value determinations on rubber samples (i.e. 100% amorphous samples, providing a good sample/sample holder contact) resulted in a Tg(onset)-value precision of • 0.5~ and a repeatability of • l~ for this method. The reproducibility of this method was determined as • 4~ during a round robin test with seven samples, measured in twenty-three laboratories [5]. These values might increase,

14

however, for Tg-value determinations on semi-crystalline and crosslinked polymers. The sample/sample holder contact is less good for the more brittle semi-crystalline polymers while crosslinked polymers show a clearly smaller 'step-wise' heat flow increase effect compared with rubbery samples.

The disadvantages of using the Tg(onset)-value as Tg-value are discussed by Richardson [2]. Determination of the Tg-value using the enthalpy/temperature curve results in a theoretically better defined Tg-value. Software to follow this procedure is commercially available at present. In the (european) industry, however, the Tg(onset)-value method is used almost exclusively because it is not only convenient, but also yields an indication for the maximum application temperature of a polymer.

1.1.4 Meltina/recrvstallisation temperature determinations Semi-crystalline polymers generally-melt over a wide temperature range. This behaviour is related to imperfections in the crystallites and non-uniformity in their size- the smaller and/or less perfectly formed crystallites will melt at lower temperatures. The endothermic fusion effect as measured by the DSC is in many cases indicated by the temperature of the maximum heat flow (the Tm-value) and by the total heat involved in the fusion process (the Hf-value). Often reported is also the Te-value i.e. the temperature at which the last crystallite has fused.

Figure 1.3 illustrates the sensitivity of the measured Tm- value for the sample weight. The maximum sample weights possible to measure a sample weight independent Tm-value are, of course, heating rate dependent. 20 ~ sample weight & 4 mg., i0 ~ sample weight & 6 mg., 5 ~ sample weight & 8 mg., (Perkin Elmer DSC-2/DSC-7, standard aluminum sample pans). The Tm-values reported in this book have been measured on 4 mg. samples at a heating rate of 20~ unless other conditions are mentioned.

A fused sample is often subsequently cooled, to follow the recrystallisation from the melt. The resulting exothermic recrystallisation effect is usually described by the temperature of the minimum heat flow (the Tc-value) and by the total heat effect involved (Hc-value). Some advance knowledge is necessary, however, to arrive at reproducible data. Incompletely fused crystal residues remain present when the temperature of the fused sample has been too low. These residues cause the nucleation process to start at higher temperatures than would normally be the case, resulting in higher Tc-values.

Samples of a commercial polypropylene (PP) grade were heated at a rate of 20~ up to a temperature Tmax. Subsequently, the samples were cooled and reheated again. The resulting melting/recrystallisation/melting values are listed

6.0

5.5 3.1 mg

4.2 mg 5 . 2 mg

s. o fl !11 I~ \ e . l mo

4.5- \'t 12.4 mg

"~ 4.0- / ,

o~ 3.5- / 1 .,,.,,.

LL i

~ 3.o- i ~ /Z

2 . 5 -

2 . 0

1 .5

1 .0

0 . 5

6 .0

5 .5

5 .0

4 .5

4 ,0

2 .5

2 .0

1.5

1.0

o. o t ~ ~ - - - - - - - - r - - - - - - - - - r - - - - - - - T - - - - - - - 200. 0 210. 0 220. 0 2~0. 0 240. 0 250. 0 260. 0 270. 0

Figure 1.3 Temperature (~

The influence of the sample size on the Tm-value of a linear polymeric system during heating in the DSC at a rate of 20~

0.0

16

in Table 1.2. These data show that heating up to at least 210~ is necessary to avoid the so-called self-seeding effect. Therefore, PP samples are heated up to 220~ as a standard procedure, before recrystallisation measurements are performed.

Table 1.2 Melting/recrystallisation data of PP samples after heating the samples up to different temperatures

i, heating Tml, Hfl, ~ J/g

162.5 I00

162.1 102

162.5 97

162.5 99

162.4 88

162.2 99 I

, , ' , " l ' ,

" l ~ a x .

oc

230

220

210

2OO

190

180

2, cool ing Tc, Hc, ~ J/g

108.6 i01

108.7 99

108.7 96

109.2 102 . . . .

109.3 98

110.0 98 �9 ........

3, heating Tm2, Hf2, ~ J/g

160.9 95

160.5 96

161.0 95

161.0 90

161.0 95

161.2 98 ,,,

a. 4 mg. powder samples b. heating/cooling rates 20~

A series of heating, cooling and heating scans is the general approach to get an impression of the melting/recrystallisation behaviour of a semi-crystalline polymer. The Tml/Hfl-values are influenced by the thermal history of the sample. The Tc/Hc-values are characteristic for the recrystallisation of the polymer under standard (thermal) conditions. Finally, the Tm2/Hf2-values can be used to compare different samples recrystallised under identical conditions.

The Tml/Hfl values listed in Table 1.2 are giving an impression of the repeatability of these measurements: Tml-value, 162.4~ • 0.2oc Hfl-value, 97 J/g • 5 J/g The base-line drawing procedure is the main reason for the relative low repeatability of the Hf-value determination. The self-seeding effect, clearly influencing the Tc-value, makes calculation of an average Tc-value meaningless. The Hc-, Tm2- and Hf2-values are hardly influenced, thus- Hc -value, 99 J/g • 2 J/g Tm2-value, 160.9~ • 0.2~ Hf2-value, 95 J/g • 3 J/g The slightly improved repeatability of the Hc- and Hf2-values in comparison with that of the Hfl-value might be caused by the improved thermal contact between sample and sample holder after the fusion process. Nakamura [5] reports a reproducibility of • 3~ for the Tm/Tc determination. The difference between the repeatability and the reproducibility values of the Tm/Tc determinations is thus considerably higher than those found for the Tg(onset)-value determination.

17

1.2 Tg-values of car-tyre rubber systems

1.2.1 Introduction The Tg-value is an important property for tyre tread rubbers [6]. It determines to a large extent the abrasion resistance, the road holding behaviour on wet roads (wet grip), the rolling resistance and the low temperature performance. A rubber with a relative high Tg-value (about -40~ generally results in a high wet grip but also in a reduced abrasion resistance and winter performance. Moreover, the rolling resistance is high! A rubber with a relative low Tg-value (about -90"C) is giving a high abrasion resistance, a good winter performance and a low rolling resistance but a reduced wet grip. Hence, the tyre tread rubber used is often a blend of different rubbers (and sometimes oil) to obtain a compromise between the properties mentioned and, of course, the cost of the tyre.

The synthetic rubbers most frequently used for car tyres are emulsion and solution styreen/butadiene random copolymers (ESBR and SSBR) and butadiene rubber (BR). Truck tyres, however, often contain a certain amount of natural rubber (NR) or its synthetic version isoprene rubber (IR). The Tg-value of BR rubber as such can vary, depending on its chemical structure between -100~ and -20~ the Tg-value of SBR can, in principle, vary between -100~ and 100~

1.2.2 The Ta-value of BR and SSBR rubbers Butyllithium-initiated homopolymerisation of butadiene results in a BR polymer containing random distributed cis-l,4, trans- 1,4 and 1,2-BR or vinyl-BR units. The concentration of the catalyst modifier and the polymerisation temperature (between 40~ and 75~ determine the concentrations of the three different components. Thus, BR rubber is in fact nearly always a terpolymer and its Tg-value can be described by means of the Gordon-Taylor relation [7]. This relation is written in its general form as:

Wi.Ai. (Tg - Tg, i) - 0 (1.1)

where: Wi = the weight fraction of monomer i, Ai = a constant characteristic for monomer i, Tg - the Tg-value of the co- or terpolymer, Tg, i = the Tg-value of the homopolymer of monomer

i; by convention Tg, i+l > Tg, i.

Constant Ai represents the difference in the specific thermal expansivities, AEi, above and below the Tg of the homopolymer of monomer i. This equation can be written for BR in the following form which is explicit for Tg-

Tg(BR) = Wc.Tg.c + W~.Ki,Tg.t + Wv,K2.Tg,v Wc + Wt.KI + Wv.K2

(1.2)

18

where- Tg(BR) = the Tg-value (Kelvin) of BR terpolymer, Wc,t,v - the weight fraction of cis-l,4, trans-l,4

and vinyl BR (I,2-BR), Tg,c,t,v - the Tg-value (Kelvin) of the 100% cis-BR,

trans-BR and vinyl-BR homopolymers, Kn = AK(n+I)/A~I.

The experimental values available for the constant Kn, in general, do not agree with those predicted by the considerations of Gordon and Taylor. Wood [7] suggested, therefore, to consider Kn as a characteristic parameter for the particular copolymeric system, not necessarily related to the A~ values of the homopolymers.

Ghysels et al. [8] used ten lithium catalysed BR samples with widely differing compositions, to calculate the (DSC) homopolymer Tg-values and the constants K1 and K2-

Tg, c = 164 K (-I09~ K1 = 0.75 Tg, t = 179 K (- 94~ K2 = 0.50 Tg,v -- 257 K (- 16~

Introducing these values in equation 1.2 is giving"

Tg(BR) = 164.Wc + 134,W~ + 129,Wy wc + 0.75wt + 0.50wv

(1.3)

where- Wc + Wt + Wv = 1.0

The difference between the measured and calculated Tg-values of the ten BR samples proved to be <_ 0.5~ This difference increased to maximally 2~ by including the Tg-values of five commercial (Co-, Ni- and Ti-catalyst based) BR systems.

To calculate the Tg-value of SSBR, equation 1.2 was extended to:

Tg(SSBR) = 164.Wc + 134.Wt + 129.Wv + K3.Ws.Tu.s WC + 0.75Wt + 0.50Wv + K3.Ws

(1.4)

where: Tg, s = the Tg-value of polystyrene i.e. 378 K, Ws = the weight fraction of styrene monomer.

The Tg-values of six SSBR samples measured were used to calculate an average value for K3. The obtained value of 0.6 was subsequently substituted into equation 1.4 resulting into-

Tg(SSBR) = 164.Wc + 134.Wt + 129,Wv + 227.W~ Wc + 0.75Wt + 0.50Wv + 0.60Ws

(1.5)

where- Wc + Wt + Wv + Ws - 1.0

The measured and the calculated Tg-values are listed in Table 1.3. The average value of Tg(experimental) - Tg(calculated) is -2~ + 4~

19

Table 1.3 Composition and Tg-values of SSBR samples

sample code

B 473

B 476

B 475

EI66AC

B 4193 ,

SSBRI

sample composition cis trans vinyl styrene

0.364 0.427 0.073 0.136

0.341 0.450 0.073 0.136

0.183 0.202 0.375 0.240 ,,

0.147 0.272 0.377 0.204 , ,

0.099 0.112 0.534 0.255 . . . . . . . .

0.297 0.465 0.085 0.153 SSBR2 0.135 0.270 0.405 0.190 BRI 0.970 0.020 0.010 0.000

, , �9 .... ,, ~, , ,,, , ,

Tg(e)

192

193

Tg(c)

242

229

266

197 232 163

196.0 ,

196.5

240.1

234.9

262.0

202.0 235.0 164.7

,

a. IR composition data, b. B473 to B4193 experimental systems, c. SSBRI and SSBR2 are commercial SSBR grades, d. BRI is a commercial BR rubber grade.

Tg(e) - Tg(c)

-4.0

-3.5

+1.9

-5.9

+4.0

-5.0 -3.0 -1.7

, ,

Equation 1.5 permits us to give an impression of the Tg-value of SSBR, as a function of the vinyl/styrene content, assuming Wt = 2Wc, see Figure 1.4. The vinyl content is usually expressed as the fraction of the BR part whereas the styrene content is expressed as the percentage on the total (SSBR) polymer. This notation is also used in Figure 1.4.

1.2.3 Th~ Tg-value after blending and oil-~tention of BR and SBR rubbers The vulcanisate properties of SSBR car tyre rubbers are often adjusted by blending with other rubbers e.g. natural rubber, BR rubbers or high styrene ESBR rubber. Usually, these rubbers are not compatible [9]. Figure 1.5 shows, for example, the DSC thermograms of SSBR/BR (75/25) samples mechanically blended in an internal mixer at 50~ and solution (cyclo-hexane) blended. The Tg-value of the BR phase (-II0~ is unaffected; however, the glass-rubber transition of the SSBR phase is influenced on the low temperature side. The two transition effects clearly present are an indication for the non-compatibility of these two polymers.

Blends of low (8.5 %wt) vinyl SSBR and high vinyl (40.5 %wt) SSBR, in spite of their relatively large difference in vinyl content, almost fully miscible at room temperature. This is indicated by the occurrence of only one glass-rubber transition temperature effect for both the mechanical and the solution blended systems, see Figure 1.6. Only the temperature shift of the Tg-value and the increased transition width (30/40oC in comparison with about 20~ for both pure polymers) are an indication that this system is a blend of two polymers. The slightly different curves might reflect the different blending techniques used. Figure 1.7 shows the DSC Tg-values

20

Figure 1.4 Tg/1.2 BR-styrene relation (SSBR rubber)

2 0

10

0

C) - 1 0 d~ r - 2 0 -1:3

6 - 3 0

I - 4 0

I/) c -50 0

I- - 6 0

0 u) -70 d:]

- 8 0

- 9 0

- 1 0 0 ! ~ , 0 .00

( 7 8 ) . . . . . . . ~/~(88) / /-+ ( 5 0 )

4 0 )

(B)

( ) ~ w t . =f, y r e n e I ( o n t o t a l p o l y m e r ) (N o f t h e t o t a l . BR par ' )

0 .20 0 .40 0 .60 0 .80 1.00

1.2 BR f ract ion

21

O. '=JO

O. 2S

O) 0.20

3= 0

U. 0.15

O I

0. I0

O. OS

i n t e r n a l m i x e r ( 5 0 ~ ~ - - ~ . ~ e p a r e d s a m p l e

/ / /" " ~ ~ , / ~ ~ ~ / ~ cyclohexane solution

O. O0 ' ' " . . . . . . . I " - [ - i ' " -120. 0 --[ tO. 0 - tO0.0 "gO. 0 -BO. 0 -70. 0

Figure 1.5 Temperature (~ Tg effects of SSBN/BR (75/25) blends

I I - 6 0 . 0 -SO.O

-0. 3o

-0, 25

-0. 20

�9 0. IS

-0. tO

-0.05

O. O0 - 4 0 . 0

0.24

O. 22

O. 20

O) O. 18

~ " O. 16

,00 0 . |4 LL "~ 0.12 O)

I o. lo O. 08

O. 06

O. Oa

0.02 -

O. O0

sample prepared from a_

. , / / intemal mixer (50~ " ~" / / prepared sample

/

-0.2a

"0. 22

- 0 . 20

"0. I8

"0. 16

-0. ia

-0. IZ

"0. I0

-0. 08

-0.06

-0. O,t

I I I ...... I ' I ~' i . . . . . . I

-90.0 --3~. 3 -70.0 -50.0 -~0. 0 --~'0.0 -30.0

Figure 1.6 Temperature (~ Tg effects of low vinyl content SSBR / high vinyl content SSBR (75125) blends

"0. 02

"0. O0

22

of two solution blended low vinyl content SSBR/high vinyl content samples plotted as a function of the high vinyl content SSBR weight fraction. This figure also demonstrates that the Tg-values of these blends can be calculated with the aid of the Gordon-Taylor relation-

Tg(blend) = Wl~Tgl + W2.KI.Tg2 (1.6) Wl + K1 .W2

where: Tgl = the Tg-value low vinyl content SSBR (197 K), Tg2 - the Tg-value high vinyl content SSBR (232 K), WI,W2 = weight fraction of both SSBR samples, K1 = system constant.

Introducing the measured Tg-values of 200.5 K for the low vinyl content SSBR/high vinyl content SSBR 75/25 (solution) blend and 217.5 K for the 25/75 (solution) blend of the same polymers resulted in an average K1 value of 0.41. Using this value for KI, equation 1.6 can be written as:

Tg(low/high vinyl SSBR blend, K) = 197.WI + 95.W2 W1 + 0.41W2

where" W1 = weight fraction low vinyl content SSBR, W2 = weight fraction high vinyl content SSBR, W1 + W2 - 1.0

(1.7)

The good agreement in figure 1.7 between the calculated Tg- values and measured Tg-values (within 1.5"C), indicates the usefulness of equation 1.7.

Oil is also often a component of the car tyre rubber compound. It is blended with the pure rubber forming the so-called oil- extended rubber phase. Usually an aromatic oil is used; such an oil showed a Tg(onset)-value at 232K (-41~ But also naphtenic oil with a Tg(onset)-value of 208 K (-65~ is used.

A few experimental data were available based on SSBR (Tg-value = 197 K, three systems) and on ESBR (Tg-value = 215 K, two systems) both extended with an aromatic oil. The Tg-value of the rubber phase is in both cases lower than that of the oil phase i.e. the SSBR/ESBR rubber phases are for Tg-rubber < temperatures < Tg-oil, 'filled' with glassy oil 'particles', resulting in an increased Tg-value of the rubber after the oil addition. The experimental values could be fitted again satisfactorily using the Gordon-Taylor relations-

Tg-value (SSBR/aromatic oil, K) = 197.W1 + 123.W2 W1 + 0.53W2

Tg-value (ESBR/aromatic oil, K) = 215,WI + 91.W2 W1 + 0.39W2

(~.8)

(1.9)

2 3

Figure 1.7 Tg value of high and low vinyl content SSBR blends (solution blended SSBR systems)

+ calculated A measured

- 4 0

- 4 5

- 5 0

= - 5 5 r @

I:1:: m - 6 0 Or~

@

> e, - 6 5

I - -

- 7 0

- 7 5

- 8 0 _ , I , ,,I . . . . . . I I, , I t I i , , , I

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 high vinyl content SSBR weight fraction

1.O

24

where- W1 = weight fraction SSBR or ESBR rubber, W2 = weight fraction of aromatic oil. W1 + W2 = 1.0

One experimental value was available for a SSBR sample (Tg- value = 232 K) extended with a naphtenic (Tg-value = 208 K). The Tg-value of the rubber phase is now higher than that of the oil phase i.e. the oil will act as a plasticiser in the temperature region between Tg-oil and Tg-rubber and the Tg- value of the rubber is ~ecrease~ after the oil addition. Assuming that also in this case the experimental values are described by the Gordon-Taylor relation, the equation might hold:

Tg-value (SSBR/naphtenic oil, K) = 208.WI + 362.W2 W1 + 1.56W2

(1.10)

where- W1 weight fraction naphtenic oil, W2 weight fraction SSBR. W1 + W2 = 1.0

Figure 1.8 shows that the agreement between the experimental data and the calculated values is satisfactory.

25

- 4 0 ;BR

T g - v a l u e of o i l - e x t e n d e d SSBR and ESBR sys tems

& m e s s .

~ / ~ $

aromatic oil

- -45

C3 ,: - 5 0 @

D - 5 5 -a

r -a

r �9 ~ - 6 0 X r I

o . . . .

o - 6 5

D . . . . .

> I - 7 0 C~ F-

ESBR

- 7 5 SSBR

- 8 0 _ _J_

0.0 0.1 I ~. I , ~ I , , I _

0.2 0.3 0.4 0,5

n a p h t e n i c oi l

. ! . . . . . . . . . . . . . . i

0.6 0.7

!

0.8

. , | , �9

0.9 1.0

Figure 1.8 Oil, we igh t f rac t ion

26

1.3 Recrystallisation and fusion of polypropylene

1.3.1 IntroductiQn Three different polypropylene (PP) modifications can be distinguished- the atactic, the syndiotactic and the isotactic modification. The atactic modification is an amorphous polymer with a Tg(onset)-value of -21~ The syndiotactic modification, made with a stereospecific homogeneous metallocene catalyst, is a semi-crystalline polymer (crystallinity about 25 %wt.) with a Tm-value of about 133~ [I0]. The isotactic modification, made with a stereospecific heterogeneous Ziegler Natta catalyst is also a semi- crystalline polymer (crystallinity about 50 %wt.) with a Tm- value of about 160~ and contains nearly always 2 %wt. - 5 %wt. of atactic material.

At present, isotactic polypropylene (i-PP) is commercially by far the most important system of the three modifications mentioned above. During crystallisation from the melt, i-PP is usually in the u form, which has a monoclinic crystal lattice with a Tm-value of about 160~ The occurrence of a ~ form (with a hexagonal lattice and a Tm-value of about 152~ is also possible during crystallisation under stress. Besides, a third (gamma) form with a triclinic crystal lattice is possible under exceptional circumstances [II].

1.3.2 Additives acting as nucleating agents for PP The recrystallisation from the melt of standard, commercial i- PP grades, is characterised by a Tc-value (see 1.1.4) of about II0~ i.e. about 50~ undercooling is necessary for spontaneous recrystallisation. Such an amount of undercooling causes relative long duty cycles during injection moulding processing. Nucleating additives are, therefore, often used in the industrial practice to decrease the injection moulding cycle times or to improve optical/mechanical properties by reducing spherulite sizes. The most efficient nucleating additives for PP, like 4-Biphenyl carboxylic acid and 2- Naphtoic acid are able to increase the Tc-value from about ll0~ to about 130~ [12].

Talc is often used as nucleating agent but also carbon black or glass fibres, added for other reasons, can act as nucleating agents. This is shown by the results of a series of Tc-values measured on PP samples with talc, carbon black and talc + carbon black, listed in Table 1.4. The same Tc-values are plotted as a function of the total additive content in Figure 1.9. These data indicate that the Tc increasing effect of both additives seems to be the same i.e. their contributions can be added up. The Tc-value of 125~ seems to be the maximum value reached due to the addition of about 1 %wt. of talc, carbon black or a combination of these two additives. The obtained effect nearly disappears however, if a considerable higher amount (i0 %wt.) of talc is added. The increase in Tc-value of 15~ is accompanied by a Tm2-value

2?

Figure 1.9 PP Tc-value/additive content relation (additives: talc/carbon black)

128

126

124

122 O

120

> I 1 1 8 O F- & 1 1 6 el.

114

112

110

&

_

~ p p Tm2-value/Tc-value J additives: talc/carbon black

166

~1~ 165

154 A

~i~ ororo ~ 163

~E 161

I--- 160 n n 159

15e

157

,,, A

156 �9 , - . . . . . . . t 0 8 112 11(~ 120 124 12"

Tc-value, ~

1 0 8 ! 0 0

I . . . . . I I . . . . I I J J . . . . I

0.2 0.4 0.6 0.8 1.0 1.2 1 4 1.6 1,8

Addi t ive content, %wt.

28

increase of 6~ i.e. the amount of undercooling necessary for recrystallisation decreases from 48~ to 39~

Table 1.4 Effect of talc and carbon black addition on the Tc- v a l u e o f PP .

~, ,', , , J ~ , " , , I , , , ' , , " " , , ,, ' l , | , , ' , , , , , , , , ,

talc i carbon total 1 Tc- Hc- Tm2- c o n t e n t , b l a c k a d d i t i v e .

content, content, value, value, value, % wt. i % wt. i % wt. ~ J/~ ~

,,, , ~ ...... I . . . . I . . . . .

0 .00 ! 0 .00 l 0 .00 ii0 91 158 . . . . . . , , , I , , I I I

o . o 5 , o . o o 0 . 0 5 , , 9 7

0.19 0.00 , 0.19 118 95 161 , , i , , . . . . . i_ !

:

�9 0.00 0.47 0.47 120 , 91 162 �9 . . I . . . . . I , , , I , , , ! , - -

0.53 0.00 0.53 118 97 161 i i . . . . . m I . . . . . I ,= , I , ! ,

,. 0 . . , 3 8 , ' 0 . 5 9 , . 0 .97 r 124 . I 9.5 . 163

0.00 I 1.37 1.37 125 97 163

i [ . . . . . . . .

, 0.76 ~ 0.85 ~.6~ , ~25 , 96 . ~64

L ,, 9.57 1.02 I i0.59 114 95 160 , , .... ..... : . . . . . . . . . . . : . . . . L,

a . a c o m m e r c i a l PP g r a d e , b. talc added via a I0 %wt. masterbatch, c . c a r b o n b l a c k a d d e d v i a a 40 %wt. m a s t e r b a t c h , d . t h e t a l c a n d c a r b o n b l a c k c o n c e n t r a t i o n s m e a s u r e d b y TGA

analysis on the 4 mg. DSC samples.

The equilibrium melting temperature for the ~ form of i-PP is still uncertain; values in the range between 185.2~ and 208.2~ are reported [13]. Whichever of these two extreme values might be the right one, the rather big difference between the melting temperature and the crystallisation temperature means that the crystalline phase of PP is very sensitive for its thermal history i.e. for annealing processes.

1.3.3 Annealinu experiments with i-PP. A series of annealing experiments was performed using an experimental, PP powder sample. This high isotactic content (96 • 1%wt. isotactic triads) i-PP had a molecular weight (Mw-value) of about 300.000 and a Mw/Mn value of 5.0, it was coded HH-SB-35.

The DSC samples (4 mg.) were heated up to the annealing temperature T(a) and stored there during a certain time t(a). Subsequently, the samples were cooled down to 20~ and

29

reheated up to 220~ The heating/cooling rate was 20~ minute.

Fillon et al. [14] showed already that the upper side of the endothermic fusion maximum forms the temperature region in which PP is most sensitive for annealing. A few scouting experiments resulted indeed in a strong Tm- and Hf-value increase for T(a) = 163~ A series of experiments with t(a) values between 5 and 60 minutes at T(a) = 163~ learnt that a t (a) value of at least 30 minutes is necessary to reach an equilibrium situation, see Figure i.i0. The curves in this figure clearly illustrate that this time is necessary to convert the less perfect crystalline fraction originally present (standard Tm-value about 162~ into a more perfect state as shown by the increase of the Tm-value from 162~ to about 175~

Subsequently, a series of experiments was performed with T(a) values ranging from 146~ up to 167~ while t(a) was kept constant at 30 minutes. The Tm-value of this PP sample proved to increase from 161~ to 176~ due to annealing between 146~ and 163~ see Table 1.5. The shape of the fusion curve starts to change considerably for annealing temperatures > 163~ see Figure I.Ii. The perfection of the whole crystalline fraction improves due to annealing at temperatures up to 163~ At T(a) values of 164~ or higher a lesser part of the crystalline fraction can improve still further. This 'high Tm-value fraction' disappeared at a T(a) of 167~ While this 'high Tm- value fraction disappeared, the crystal fraction with the 'standard' Tm-value of about 162~ increased again, see Figure 1.12.

In view of the results shown in Table 1.5 and Figure I.ii, four different annealing regions can be distinguished in the 'standard' fusion curve of this i-PP sample represented in Figure i. 13 :

Annealing region I, T(a) < 150oC The fusion curve becomes asymmetrical on the low temperature side; slightly higher Tm- and Hf-values. Annealing region II, 150"C _< T(a) _< 163oC The fusion curve is more or less symmetrical, the Tm-values increase with T(a) and the Hf-values are going through a maximum. Annealing region III, 163"C < T(a) < 167oC Fusion curve with two maxima, the Tin-value reaches its maximum value (179.5~ the Hf-value becomes zero, this is accompanied by- an increase of the Hf'-value up to 100 J/g, while Tm' is constant (about 164oc), see Figure 1.12. Annealing region IV, T (a} > 167 "C More or less symmetrical fusion curve, Tm'- and Hf'-values decrease to the 'standard' values with increasing T(a) values.

7 . 0 - ' ' "

6 . 5 -

6 . 0 -

5 . 5 -

5 . 0 -

~ 4 . 5 -

4.0- o ~ 3 . 5 -

I 3 . 0 -

2 . 5 -

2 . 0 -

t . 5 -

1 . 0 -

0 . 5 -

~ Rnnealed a t , 163~ d u r i n g " / / -'~

I

R e ~ e e r e n c e po~der' mample / ! HH-SB-35 (end samp ] e)

15 mln.

5 mtn.

~ ~

" I ' f " I ' t ' ' i ' 130.0 140.0 t50.0 160.0 170.0 t80.0 t90.0 T e m p e r a t u r e (~

Figure 1.10 DSC fusion endotherms of reference sample HH-SB-35 measured after annealing at 163~ using different annealing times

.0-

6 .5-

5.0-

5 ,5 -

...~ 5 . 0 - o'J

~.. 4 . 5 -

4 .0 - I,I,

~ 3 5 - ~) �9 "I"

3 ,0-

2.5

Re~e Pence powder s amp I e HH-SB-35 (end 8amp 1 e )

/ e

I !

/ / t /

2o-1 _~~ - ~ ~ ~ ~ J - j

t . 5 Q ~ m Im~ml l~P

1.0

30 m! nul;es anneal ed at. z

163"C

164 *C

IG5~

0 .5 - ~ , ~ [ t30.0 i40 .0 i50.0 i60 .0 t70.0 180.0

Temperature (~ Figure 1.11 DSC fusion endotherms of reference sample HH-SB-35 measured after 30 minutes annealing at different temperatures

[90.0

32

These results were successfully applied to estimate the maximum temperatures 'seen' by the product at different locations in a PP reactor system during a series of reactor trials.

Table 1.5 Results of annealing experiments on i-PP, samples 30 minutes at T(a)

!

T(a) value,

oc ..

146 ._

150

152

156 _

159

161

162

163

164

164.5

165 ,

166

166.5

167

Tc value,

oc

141.4

140.8

140.9 ,,

138.9

138.3

135.8 ,, ,

Tm value,

oc ,

161.0

163.4

165.1

168.7 .,,

172.0

174.4

175.8

175.8

178.3

178.7

179.2

179.2 ,,

179.5

,,, . ,.!

Hf value, J/g

76

81

79

78

91

109

Ii0

107

3 1

1 2

ii

2

Tm' value,

~

, 0 ,

i

161.3

163.5

164.2

164.2

164.6

163.6 ..

I Hf'

value, J/g

. . . .

- i

I

L

I

L

69 , i

81

89

97

99

i00 i

- ,

* standard procedure i.e. heating to 220~ followed by cooling to 20~ and reheating at a rate of 20~ resulted in- Tc-value = I13~ Tm'-value = 159~ Hf'-value = 85 J/g.

Tm-v

alue

, ~

@

O~

O~

@

O~

"q

"q

"q

"q

".q

CD

-~

0 i~

-I~

C~

CO

0

PO

.1~

~ ~

0 -1~

> ._

~

~ --4

~ co

(3

0"I

0 PO

~ @

CD

-~

-~

-~

-~

-~

0 0

0 0

0 Po

~ O~

Oo

0 0

0 0

0

Hf-v

alue

, J/

g

.L. r

I=

L,,

AE

'ro

0 �9

-,..

~ 0 -.g

-g

0 r ~

241 2 2

2.0 -I

1.8

~ i . 6 v

o J.4 u. =i i .E (D -I-

i.O

PP a n n e a ! ! n g d u r i n g 3(] mlnut, e8 a t Ta (deg . C)

r e g i o n t . Ta < t58 C r e g i o n 2. tS;] C .< Ta .< t63 C r e g i o n 3. 163 C < Ta < 187 C r e g i o n 4. Ta I t67

0.8

0.6

0.4

0.~ 1 . .~ ,o, ,..__,o,,~, o~ ~._~ ~' _ t ' - ~ 1 7 6 r �9 0.0

too . o " t 2 0 . o " t4o . o t 6 o . o t e o . o

Figure 1.13 Temperature (~ The "standard" melting endotherm of reference sample HH-SB-35 with the four different annealing regions indicated

35

Fillon et al. [12] proposed an efficiency factor for the evaluation of nucleating additives for polymers-

nucleating efficiency (NE) = I00. (Tc.na -Tcl) (Tc2,max. - Tcl)

(1.11)

where- Tc,na = Tc-value of the system with the nucleating

additive i.e. 125@C for the PP system with 1.05 %wt. talc and carbon black (Table 1.4),

Tcl = Tc-value of the reference system i.e. II0~ for the system in Table 1.4,

Tc2,max. = Tc-value of the system self-nucleated to saturation [12].

Tc2,max. was not measured for the system in Table 1.4. The Tc- value of 141.4~ given in Table 1.5 can be used, however, as a reasonable approximation (both PP systems are made with comparable catalyst systems). Using this Tc2,max. value results in a NE-value of-

NE = I00. (125 - 110)/(141.4 - ii0) - 47.7 i.e. 48%

Fillon et al. [12] report a NE-value of 32% for PP with 1%wt. of talc as nucleating additive. The difference between both NE values is not too bad considering the Tc2,max approximation, the differences in the used experimental methods and the different PP systems investigated.

36

1.4 Side-chain crystallisation in poly(l-olefin) s

1,4,1 Introduction The presence of long (linear) side chains in branched polymers can cause side chain crystallisation. It is important to distinguish main chain from side chain crystallisation for the effect of this difference on the product properties can be considerable.

The Tm- and Tc-values of a polymeric system increase in general as a function of the chainlength i.e. the system's molecular weight. Both values become constant at higher molecular weights or go through a kind of maximum value [2]. This makes discrimination between main chain and side chain crystallisation on basis of the side chain length easy" - Side chains are in comparison with the main chains usually

short and the Tm- and Tc-values of side chain crystallisation effects will, therefore, in general increase with increasing side chain length.

- The presence of side chains hampers usually the main chain crystallisation. Increasing side chain length will result,

in general, in decreasing main chain crystallisation Tm- and Tc-values.

A series of poly(l-olefin) s was analysed by DSC, offering the possibility to map the differences between side and main chain crystallisation.

1.4.2 Crystallisation in poly(l-olefin)s A series of C6 up to C18 ~-olefin fractions prepared by the so-called SHOP process were polymerised with a Ziegler-Natta catalyst system. The purity of these fractions was > 98 %wt. The (peak) molecular weights of the polymers proved to be > 200.000; NMR analysis showed atactic material and some stereo- regularity. The results of the fusion/recrystallisation measurements (see 1.1.4) are listed in Table 1.6.

Table 1.6 Results of DSC analysis of poly(l-olefin) s

Cn frac-

tion . . . . .

C6 m

C8

.~ C10

1212 I

~, _ ,

i

i,. c l ,4

i c16 L

system A. or s-C.

A

A

A

s-C

s-C .

s-C

s-C , ,

Tg- value

oc

-47 . ,

-73

-75

T

Tc- value

oc ,

-67

-22

8 . . . .

24

36 | .,,

, . w . �9 z

Hc - Tm2 - I value J/g

15

38

68 . . . . .

79

95

* A. = amorphous; s-C. = semi-crystalline

value ! ~

i

. . . . .

22 i i

35

46

58

66 �9 ... ,% ,

3 . 5 0 -

3 . O0 -

,!

2 . 5 0 -

.9o LL �9 ,-, 2 . 0 0 - (9

::E

1 . 5 0 -

1. O0

0 . 5 0

O. O0

- 1 0 0 . 0

RECRYSTRLLISRTION FROM THE MELT OF SHOP LINERR RLPHR-OLEFINE BRSED POLYMERS

. . . . . . . . . . . . . . --..-.C~.I~,~SHOP C 1 2 - S H O P C ! 4 - S H O P C I ( ] - S H O P

~ _ / "-~.,.. " ~ . ~ �9

Temperature (~ . . . . . . . . . . . . . . . ~ ............................. t . . . . . . . . . . . . . . . 1 ................................. I . . . . . . . . . . . . . . . .

- 7 5 . 0 - 5 0 . 0 - 2 5 . 0 O. 0 I . . . . . . . . . . I -

2 5 . 0 5 0 . 0

4 . 0 0 - = . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . - . . . . . . . . . . . . .

Figure 1.14

75.0

L~ ,.J

38

The C6 based polymers (side-chain: -[CH2]n-CH3, n=3) and the C8 based polymers are amorphous systems which did not crystallise even at low temperatures. The CI0 based system is 'as received' amorphous but crystallises at low temperatures. The C12 to C18 based polymers are semi-crystalline systems. Figure 1.14 shows the recrystallisation exotherms of the CI0 to C18 systems. These curves show a strong increase of the extent of the crystalline phases with the chain length. This increase is somewhat distorted by expressing it in J/g. But, if expressed in J/mol, the difference in Hc-value between the Cl0 based polymer (Hc = 0.ii J/mol) and the C18 based polymer (Hc = 0.38 J/mol) is still clearly present!

The Tc- and Tm2-values of these poly(l-olefin)s are plotted in Figure 1.15 as a function of the number of C-atoms in the side-chain. Besides, Tm- and Tc-values of the main-chain crystalline phases as present in polypropylene (PP), poly l- butene (PIB) and poly l-pentene (PIP):

PP : Tm = 162~ Tc = II0~ PIB: Tm = 124~ Tc = 67~ PIP: Tm = 70~ [3],

are also plotted as a function of the number of C-atoms in the side-chains. The differences earlier decribed between the two types of crystalline phases are clearly present.

0 0

d .,1,.

> 6 I:::: E F-

2 4 0

2 0 0

160

120

80

4 0

0

- 4 0

- 8 0

3 9

Figure 1.15 Tm/Tc-va lues of po ly(1-o le f ins)

+ Tm A Tm s.c. m.c.

o T c + /T/. C.

Tc S.C.

0

\ \ \ \ \\ \

0

| ..I

2

d

> r

Hc-value poly(1 -olefin)s

,oo - y

8 O

7 0

6 O

6 O

4 O

310

1 0

0 45 8 1 0 1 2 14 1 6 1 0

Number of C-atoms in side chain

..i _ ~ +J J

4.

J" 4-

I ~ . . . I , I , I ..... , I ~ ....

4 6 8 10 12

f i - ' l - "

I I "l " j

�9 .-I..- j f " l " " " " 4" I "

. . . ,,.....,

| ! | I .

14 16 18

Number of C-atoms in side chain

40

1.5 Chemical reactions monitored by DSC

1,5,1 ~ntroduc~ion The cure reaction of thermosetting resin systems is the subject of many publications about monitoring chemical reactions by DSC. In much of this work is tried to derive kinetic information from the heat of reaction measured during both scanning and isother~al experiments [2, 15]. However, the results reported by Wisanrakkit and Gillham [16] illustrate the insensitivity of the DSC technique for small (residual) cure exotherms. They show that the development of the Tg-value during isothermal DSC experiments is offering a much more sensitive measure for the conversion of a curing system. Besides, the heat of reaction measurement can not be used for all thermosetting resin systems. There are thermosetting resin systems which give no exothermic effect at all during their cure.

This problem is illustrated by the results of four experiments performed with a Perkin Elmer DSC-2, using a heating rate of 20~ Figure 1.16A shows the cure exotherm of an epoxy powder coating system cured with an amine based curing agent. The exothermic effect is strong, the begin and end temperatures for a (partial) integration procedure can be defined easily. Figure 1.16B shows the cure exotherm of an epoxy powder coating system cured with a phenolic-OH based curing agent. This cure exotherm, although clearly smaller than the proceeding one, can still be treated with the same calculation techniques.

Figure 1.16C shows the exothermic effect of an epoxy powder coating system cured with an anhydride based curing agent. A certain exothermic effect seems to be present but a (partial) integration procedure is with this result impossible. Figure 1.17, finally, shows the DSC thermograms of an epoxy powder coating system also cured with an anhydride based curing agent, which shows no cure exotherm at all. The cure process of these resins consists of a number of both exothermic and endothermic reactions; the DSC measures the totall amount of heat released which thus even can be nearly zero! The Tg-value increase seen in the second scan is the only indication for the occurance of a cure reaction during the first DSC scan. These examples comfirm the conclusion of Wisanrakkit and Gillham [16] that the Tg-value development offers in many cases the best possibility to characterise the cure process of a certain thermosetting resin system. An example of such a procedure is given below (see 1.5.2).

This conclusion holds especially for the examples of relative complex cure processes shown above. Less complex reactions with pure, low molecular weight components can often succesfully be studied using the exothermic reaction effect. An example of such a procedure is also given below (see ~.5.3).

,12

5=.,0

T~ ~ ~ ~ K ~-40 _

con0 scao / z.4~O

Figure 1 . 1 7 , / | The first and second DSC / ~oo .~ heating scans on an epoxy ~ | resin based powder coating ~ | system =zo

S.T~Lp..~~~A~,'o.e~-.~-~-_3~,'K 3,,O "!

"t'E"PKRA'ruRE, K I First scan ~ s~r r ,~= T~.s . tm" ~ J

43

1.5.2 The determin@~ion of the cure conditions of a powder coatina system The Tg[vaiue development of an epoxy powder coating system was used to determine the cure time (at 180~ as a function of the curing agent concentration, necessary to reach a Tg-value of at least 100~

The Tg-value of the system before cure [Tg(o)], the Tg-values after different cure times at 180~ [Tg(t)] and the maximum Tg-value due to cure at 180~ [Tg(e)] were measured using four curing agent concentrations i.e. 13.5 phr., 17 phr., 20.5 phr. and 24 phr. A typical result of such a series of measurements is shown in the inserted figure of Figure 1.18. The results of the Tg(o)- and Tg(e)-value determinations were averaged:

Tg before cure: Tg(o) = 60.5~ + I~ Tg(maximum) : Tg(e) = I08.5~ _+ 1.5~

The conversion of the cure reaction can, based on the Tg(t)- values and using these Tg(o)- and Tg(e)-values, be expressed as :

conversion x(t) = Tg(t) - Tq(o) Tg(e) - Tg(o)

(1.12)

The conversion data of the investigated system with four different curing agent concentrations are plotted as a function of the log(cure time) at 180~ in Figure 1.18. The drawn straight lines have Rval.-values > 0.993; the slopes of these curves increase slightly from 0.49 for the system with 13.5 phr curing agent to 0.55 for the system with 24 phr of the anhydride based curing agent.

The conversion of this system has to be 0.82 or more to reach a Tg-value of at least 100~ according to equation 1.12. Using Figure 1.18, Figure 1.19 can be constructed containing the information asked by the customer about this system.

!.5.~ Reactions of model compounds studied by DSC The development of acid based curing agents, to cure a new generation of UV resistant epoxy resins, required a study of the epoxy/acid reaction with model compounds. A mono- functional, liquid epoxy resin (CARDURA E5) was used as model resin; the selected model acids are listed in Table 1.7.

The reaction of a stoichiometric amount of these twelve acids with the epoxy resin was measured during a series of non- isothermal DSC experiments. These measurements were performed with a Perkin Elmer DSC-2C using a scanning rate of 10~ minute. Each experiment consisted of three heating scans:

44

Figure 1.18 (Tg-Tgo)/(Tge-Tgo) of a powder coating system versus the cure time at 180~

-}- 13.5 /k 17 phr phr

curing agent cone.

0 20.5 4- 24 pit phr

1 . 0 0

0 .90

0 . 8 0 +

0 . 7 0 0 o) F- 1 0 . 6 0 E~ E--

0 . 5 0 0 E~ F- 1 0 . 4 0 E~ F-

0 . 3 0

0 . 2 0

0 . 1 0

4 - 1 2 o

1 l O

lOO

7 0

6 0 o

+ / + ~ 4- /

/ /4 §

/ 4- !

+

/ Tg development during cure at 180~

t

1 0 0 0 2 o o o 3 o o o 4 0 o o e, o o o

Cure time at 180~ s

0 . 0 0 I /

3 0

I I I I I I I

100

i i i i i , ,I f l

1 0 0 0 . , , I I I I I I I I

1 0 0 0 0

Cure t ime at 1 8 0 ~ s

45

1 0 0 0 -

,,,,,

0

0 ' T "

r "

E

L _ _

2 O 0

10

Epoxy resin based powder coating system (cure temperature 180~

. . . . . . . . . . . . . .

Figur Time/curing agent concentration relation necessary to reach a Tg-value of the "~. cured product of at least 100~ ~ .

I I , I . . . . . . I I _ I , , l

12 14 16 18 2 0 2 2 2 4 2 6

Curing agent concentration, phr.

46

Table 1.7 Survey of the used model compounds

,,,, , ,

Acid chemical structure sol. in liquid epoxy

, res in

hexane acid CH3 - CH2 -CH2 - CH2 -CH2 - COOH t

I i ' ' ' i i

i isobutyric acid CH3-CH-COOH i ~H3

I

pivalic acid

ii , ,

hydroxy-pival ic acid

i~ ,,

cyclohexane i carboxylic acid

ii l

l-methyl-cyclo hexanecarboxylic acid

benzoic acid

ml ,,

2 -methyl - b e n z o i c a c i d i

3 -methoxy- benzoic acid


2 -ethoxy- benzoic acid

~l . . . . . . . . . . . .

4 -ethoxy- benzoic acid

~H3 CH3 -C-COOH

~H3

CH2OH

~ COOH

/-COOH

~COOH

CH3 -COOH

0 .

CH3 - ~ COOH

oO coo.

~ -CH2 -CH3 COOH

L

CH3 - CH2 -0 O COOH

liquid epoxy resin- ~H3 ~,O

c H 3 - c - c - o - c . 2 CH3

4 '/

- the first scan from 20~ up to 250~ to determine the extent and temperature location of the reaction exotherm,

- the second scan from -120~ up to 250~ to measure the Tg- value of the reaction product and a possible residual exothermic effect,

- the third scan from -120~ up to 20~ to check if the Tg- value still increased after a second thermal treatment.

The measurements were performed with the samples in high pressure capsules (internal pressure maximally 150 bar) to avoid sample loss due to evaporation.

A possible mixing problem had to be solved first. Six of these twelve acids are soluble in the liquid epoxy resin at room temperature and the DSC samples were made using a standard masterbatch procedure. The other six acids, however, proved to be insoluble in the liquid epoxy resin. Weighting the insoluble acids directly into the DSC high pressure capsules upon the liquid epoxy phase was the first option. Subsequently two epoxy/acid systems, with one in the resin soluble acid and one insoluble acid, were measured four times to detect possible differences in the repeatability of these measurements:

soluble acid liq. epoxy resin /hexane acid,

Tmax. (exotherm) = 161 _+ 2~ dH(heat of reaction) = 87 _+ 1.5 kJ/mol

insoluble acid liq. epoxy resin Tmax. (exotherm) = 152 _+ 2~ 2-m.benzoic acid, dH(heat of reaction) = 79 + 2 kJ/mol

These results confirmed that the experimental procedure followed worked satisfactory. Figure 1.20 shows a typical result. The reaction exotherm of the epoxy/benzoic acid system is measured during the first scan. This effect is characterised by a Tmax.-value of 139~ and a dH-value of 87 kJ/mol. No sign of a residual reaction exotherm is measured during the second scan. The Tg-value increased from -I09~ for the liquid epoxy as such to -48~ after the reaction with benzoic acid. A Tg-value of -48~ was also measured during the third scan, indicating that no (detectable) further progress of the reaction occurred after the first scan. The results of the measurements performed in this way are listed in Table 1.8.

The first six systems listed in Table 1.8 show that the Tmax.(exotherm)-value of the epoxy/alifatic acid systems is related with the strength (pKa-value) of the acid used. The inductive action of the hydroxyl-group in hydroxypivalic acid, for example, causes a pKa-value decrease from 5.03 (pivalic acid) to 4.50 (hydroxypivalic acid). This difference results in a Tmax.(exotherm)-value decrease of 24~ Subsequently, a still stronger acid, di-ethylmallonic acid (DEMA), was tried.

~. 51~ . . . . . . . . . -

O E l Or)

< o

8. 25

Figure 1.20 The reaction of liquid epoxy resin / benzoic acid _ , j ~ i as measured by DSC . . . , , ' w

jr=" _ f

r

~ "

j .,,

s " i

o

J f I

B.eB - : .. , . 0 see.ca 2a.ee 2=.=

2 n d s c a n

T E M P E R A T U R E ( K ) I I , I . . . . . . I . . . . . I ~ - I

3=. ~ 34~. ~ ~e . ~8 42~. 00 468. ~0 s ~ . ~

OO

49

Table 1.8 The results of DSC model experiments with liquid epoxy resin/acid reactions�9

.... �9 . . . . . . . . . . . . ~,, .., . . . . . . . . , . . . . . [ .... ,

Acid pKa Tmax. dH Tgl value value, value, (Tg2)

~ kJ/mol ~ u | , , , , �9 _, , �9 ,

1 -methyl- cyclo hexanecarboxyl 176 76 -60 ic acid (-60)

pivalic acid 5.03 171 74 -66 (-66)

|�9 ,,, ,, , ,,, ,

cycl ohexane c a r b o x y l i c a c i d

�9 �9 ,,

h e x a n e a c i d

i:

4.90 167 82 -62 . . . . . L (,-62 )

isobutyric acid

,_

hydroxy- ,. pivalic acid

4 - ethoxy- benzoic acid

2 -ethoxy- b e n z o i c a c i d

.. ,,,


n~ ................

2-methyl- benzoic acid

i~ . . . . . . .

3-methoxy- ~benzoic acid

benzoic .... acid , - ~,,,,,,, ~,,,,,

4.88 161 87 -85 [ ( -85)

- . ..... , , �9

4.86 160 76 -79 ! ( -80)

, J . . . . . . .

4.50 ~ 147 63 -52 (-52)

4.80 181 54

4.21 174 80

4.47 173

3 . 9 1 152

_

I 4 09 141

! _ ,, |

i '

4.19

-34 (-3.4)

-41 (-41)

59 -36 (-36)

_ _. ......

79 - 5 2 (-52)

.... _- . . . . . . ,, -!

, . . . . . m ,

99 -42 , .'

139 87 -48 (-48)

,

50

185.00

Figure 1.21 The Tmax.-value versus the pKa-value (liquid epoxy resin/acid model series)

+ a l / fa t i c A a r o m a t i c a c i d s a c i d s

176.50

168.00

/k A

159.50

1 5 1 . 0 0

142.50

134.00

o

af . . . . . ,

E I,,-

.A

A /k

125.50

117.00

108.50

1 0 0 . O O r , i

2.80 3.20

pKa-value I J I

3.60 4.00 I ,. ~..

4 . 4 0

I

4.80 5.20

53.

~2H5 di-ethylmallonic acid �9 HOOC-C-C00H

~2H5

The epoxy/DEMA system resulted in a Tmax.(exotherm)-value of I03~ and a dH-value of 74 kJ/mol. The pKa-value of DEMA is about 2.9. All the measured Tmax.(exotherm)-values are plotted as a function of the pKa-value in Figure 1.21. it illustrates clearly that a relation between Tmax.(exotherm) and the pKa- value which is found for the alifatic systems does not hold for the aromatic systems.

The aromatic acid systems show other effects. The Tmax. (exotherm)-values of three of the aromatic acids show, for example, a clear steric hindrance effect. The Tmax.(exotherm)- values for benzoic acid, for 2-methyl-benzoic acid and for 2- ethoxy-benzoic acid increase from 139~ to respectively 152~ and 174~

The effect of a para-alkoxy substitution of benzoic acid is also clear; the pKa-values and Tmax.(exotherm)-values are increasing going from benzoic acid, to 4-methoxy-benzoic acid and to 4-ethoxy-benzoic acid. Besides, the dH-values of the last two mentioned acids are clearly lower than all the other dH-values. Such a difference in Tmax.(exotherm)-value is not present between benzoic acid and 3-methoxy-benzoic acid (meta- alkoxy substitution). The reason for the decreased epoxide reactivity due to para-alkoxy substitution might be the conjugated mesomeric structure which causes an extra negative charge on the carbonyl-group.

The dH-values of eight of these systems seem to be more or less constant i.e. 80 • 5 kJ/mol. The dH-values of the para- alkoxy substituted benzoic acid systems and that of hydroxypivalic acid are significantly lower. The dH-value of the meta-alkoxy substituted benzoic acid system is significantly higher (99 kJ/mol). The reasons for these differences are not clear.

The Tgl- and Tg2-values of all systems are equal indicating that all systems reacted only during the first heating scans. This does not mean, however, that the conversion was I00 % for all systems. The Tg-values are no indication for the conversion, in this case, due to the strong sensitivity of these Tg-values for small structural differences. The Tg-value can only be used in this situation to check if some residual reaction effect occurred.

52

1.6 Determination of the heat of vaporisation by DSC

i. 6.1 I~t.roductiQn The heat of vaporisation at 25~ (AHvap.25) of a solvent is used to calculate the Hildebrand Solubility Parameter (HSP) assuming that the evaporating solvent behaves like an ideal gas. The HSP, subsequently, is one of the three parameters used in the Nelson, Hernwall and Edwards system to describe and predict the solvent power [17].

The heat of vaporisation is usually measured with a completely closed calorimetric system permitting vaporisation experiments under controlled vacuum or pressure. The equipment developed for these measurements is rather complicated and scarcely available [18]. Farritor and Tao [19] used the convenient, wide-spread DSC technique for this purpose, accepting that this choice permitted heat of vaporisation measurements under atmospheric pressure only. Their Perkin Elmer DSC-IB was equipped with an open measuring cell system and could be used as such for vaporisation experiments. The DSC-2, -4 and -7 systems used at present, are equipped with semi-closed cell systems and have to be modified to perform vaporisation experiments. The DSC modification and the results of a series of heat of vaporisation measurements at 25~ are reported in this chapter.

1.6.2 DSC modification for th~ AHvaD.25 determination The DSC vaporisation determination is based on measuring the amount of heat necessary to vaporise a known amount of the substance. This substance is placed in the DSC measuring cell in a closed container and about I0 minutes is waited then to restore the equilibrium in the DSC cell. The heat of vaporisation determination is started, subsequently, by opening the sample container in the DSC cell and measuring the amount of heat necessary to evaporate the whole sample.

Stainless steel high pressure capsules (Perkin Elmer) provided with holes in the upper side of respectively 0.5 mm., 3.0 mm. and 4.0 mm. in diameter, are used as sample containers. The sample containers are closed by covering these holes with mild steel lids which can be removed magnetically. Both the sample container surfaces and the mild steel lid surfaces are polished to obtain an optimal closing action of the lids.

The aluminium cover of a DSC-2 system was replaced by a polycarbonate cover (see Figure 1.22) provided with a spring loaded, moveable magnetic system to remove the mild steel lids from the sample containers, while the DSC cell is closed. Two experimental parameters can be varied to vaporise samples with boiling temperatures ranging from 50~ to 200~ within maximal 30 minutes i.e. the sample container hole diameter and the sample weight. Besides, the DSC sensitivity can be varied.

~3

Figure 1.22 Modified DSC-2 sample cell for vaporization experiments (shown schematically)

j n fill! I' i I i ! I i l I i . . . . . . . . i I i ! ! il I j i I i I ! il I i

. i l i ~ ! i= =i . . . . . . " / , ' / / / / z l , ~

R

s t o p p ! n g b l o c k

s p r ! n g l o a d e d magnet== h o ! d e r

magnet~

mi Id s t e e l I id

r u b b e r 0 - r | n g

A: DSC cell base

B: polycarbonate cell cover

54

A series of preliminary experiments resulted in the following set of experimental conditions.

solvent boiling sample cup sample DSC point range, hole diameter, weight, sensitivity,

~ mm. mg. mcal/s

55 - 80 0.5 1 - i0 i0 80 - 170 3.0 1 - 2 I0

170 - 200 4.0 0.i - 0.5 0.I

The sample holder environment was purged with helium during these experiments, at the relative high flow-rate of 200 ml/minute. Such a flow-rate proved to be necessary to obtain flow-rate independent results. Apparently, the formation of a the vapour 'cloud' above the sample container is prevented in this way.

The standard DSC heat-flow calibration procedure with indium is not longer accurate enough, due to the change from a closed to an open cell system (the standard platinum cell lids are removed). Demineralised water is used, therefore, as calibration substance for the vaporisation experiments. The heat of vaporisation of demineralised water was measured six times which each of the sample cups. The average heat of vaporisation measured was compared with the known heat of vaporisation of water (43.9 kJ/mol., [20]) to calculate a correction factor for each sample cup-

sample cup average AHvap.25 correction hole diam., measured, kJ/mol, factor

0.5 n~n. 38.4 _+ 0.4 1.14 3.0 mm. 39.5 + 0.2 I.Ii 4.0 mm. 39.0 +_ 0.2 1.13

The isothermal (25~ experiments thus give a heat-flow versus time signal, see Figure 1.23. The time at which the mild steel lids are removed is indicated by (A) while the time that all the solvent is evaporated, is indicated by (B). Integration of this signal provides the total amount of energy necessary to vaporise the investigated sample. The base-line used for the heat of vaporisation integration is indicated by (C). The shift of the DSC base-line shown in Figure 1.23, is mainly caused by the removal of the mild steel lids. The total amount of energy measured is used then to calculate the AHvap.25 after multiplication with one of the above given correction factors.

1.6.3 Results of AHvaD.25 determinations by DSC The AHvap.25 of a series of samples with known AHvap.25 values [18] and boiling temperatures ranging from 57~ to 214~ was measured to determine the accuracy of this 'DSC' method. The results are listed in Table 1.9. Figure 1.24 shows the measured curve of n-dodecane. In spite of the (too) long

5 5

2 . 5

til. I~

I . . . . . .

I 2 3 4 5

T I M E ( m i n u t e s )

Figure 1.23 DSC curve of the vaporization of ethyl propionate

G 7 0

56

Table 1.9 The DSC AHvap.25 determination accuracy

solvent i

boiling tempera ture, ~

|

acetone 57 | . . . . . . .

methanol

ethanol

4-methyl-2 �9 pentanone

2- heptanone

n-dodecane

65

79

116

147

214 ~, . ,

AHvap. 25 exper. kJ/mol

30.8

38.0

42 .i

40.6

47.3

62.6 T. .

AHvap. 25 literat. kJ/mol

30.5

37.4

42.3

40.6

47.2

61.3

§

+I.0

+1.4

-0.5

0.0

+0.2

+2 .I

measuring time, a reasonable baseline can still be constructed. The results listed in Table 1.9 are the average values of triplicate measurements. These data show that an accuracy of + 2 % is possible with this method.

The AHvap.25 values of a series of propionate esters prepared with a new catalyst system was measured, subsequently, using this DSC method. Besides, a series of acetate esters was measured as references. The average values of the triplicate measurements are listed, together with the available literature values, in Table i. I0.

The ~Hvap.25 values of a homologous series of organic solvents increase with their molecular weight. Figure 1.25 shows that the measured values of the linear propionate samples and that of the linear acetate samples form in fact two different, linear relations (only the AHvap.25 value measured for n- pentyl acetate deviates clearly). The difference in heat of vaporisation between linear propionate and linear acetate samples with equal molecular weights seems to be disappeared for the branched systems. The influence of the structure on the heat of vaporisation is illustrated by the results of the following samples (mol. weight = 102).

0 n-propyl acetate CH3-CH2-CH2-O-~-CH3

0 ethyl propionate CH3-CH2-O-'~-CH2-CH3

i-propyl acetate CH3-~H-O-~-CH3 Ctt3

AHvap. 25 39.7 kJ/mol.

38.1 kJ/mol.

35.7 kJ/mol.

The more asymmetrical location of the -CO0- group in the acetate sample seems to increase the heat of vaporisation. Branching, however, clearly overrules this effect.

5?

0.30

0.15 0 I.U

r

0.00 ~. 0 6 12 18 24 38 36 42 48 54 68

TIME (minutes)

F i g u r e 1.24 DSC CURVE OF THE VRPORIZRTION OF N-DODECRNE RT 25 DEG. C.

58

Table I.I0 AHvap.25 values of propionate and acetate esters

solvent

, . methyl , propionate

ethyl propionate

AHvap. 25 experim. kJ/mol

�9 . . . . . .

32.5

38.1

n-propyl propionate i i-propyl propionate

n-butyl propionate .. s-butyl Propionate

.. n-Pentyl pr~pionate

43.8 41 .i

50.0 46.7

55.9

....... " '~ ........

AHvap. 25 literat. kJ /mo 1

�9 i 35.85 i

-,

43.45

39.21

,methyl acetate. i

ethyl acetate ,,

n-propyl acetate ,.~ i-propyl a~etate

n-butyl acetate s-butyl acetate

,I i-butyl acetate

n-penny I acetat e

30.0

34.6

39.7 35.7

....

45.1 41.5 42.4

53.0 . . . . . . . . ~. . ,u.~

32.29

35.60 ....

39.72 37.20

43.86

5 6

Figure 1.25 The heat of vaporisation at 25~ as a function of the molecular weight -t- n -p rop io A n - a c e O branched '4" branched

nares ra tes propion, a c e t a t e

5 3

5 0

4 7

4 4

41

3 8

3 5

32

29 I _

m

0 E

- d 0 1,0

,.,,,

c 0 �9 ==,.,

. , a , , .

- 0 C I .

> t 4== ,

0

I

59

+

0

+

Molecular weight I I , I I , 1 I

70 8 0 9 0 100 110 120 130 140 150

60

References

I. B. Wunderlich. Thermal Analysis, Academic Press Inc., New York, 1990.

2. V.B.F. Mathot" Calorimetry and Thermal Analysis of Polymers, Hanser Publishers, Munich, 1994.

3. D.W. van Krevelen- Properties of Polymers, Third edition, Elsevier, Amsterdam, 1990.

4. J. Bicerano- Prediction of Polymer Properties, Marcel Dekker Inc., 1993.

5. H. Nakamura, Thermochimica Acta, 136, (1988) , p. 163 - 178.

6. K.H. Nordsiek, Kautsch. Gummi Kunstst., 3~, (1985), p. 178- 185.

7. L.A. Wood, J. Polymer Sc., 28, (1958), p. 319. 8. A. GHijsels, Internal Shell Report, DTRS.0011.74. 9. J.E. Stamhuis, W.M. Groenewoud and J. Raadsen, Plastics &

Rubber, Processing and Application, ii, (1989) . i0. M. Antberg et.al., Macromol. Chemie, Macromol. Symp.,

48/49, (1991), p. 333 - 347. ii. E. Devaux and B. Chabert, Pol. Comm., 32, 15, (1991), p.

464. 12. B. Fillon et.al., J. of Pol. Sc.- Part B, Polymer

Physics, 31, (1993), p.1395 -1405. 13. J. Janimak et.al., Polymer, 3/, 4, (1992), p. 728. 14. See reference 12, p.1383 - 1393. 15. J.M. Barton, Adv. Polym. Sci., 72, (1985), p. Iii. 16. G. Wisanrakkit and J.K. Gillham, J. Appl. Pol. Sc., 41,

(1990), p. 2885 - 2929. 17. R.C. Nelson et al., J. Paint Techn., 42, (1970), p.636. 18. V. Majer and V. Svoboda" Enthalpies of vaporisation of

organic compounds, Blackwell Scientific Publications, Oxford, 1985.

19. R.E. Farritor and L.C. Tao, Thermochim. Acta, i, (1970), p. 297.

20. Handbook of Chemistry and Physics, 41 st edition, Cleveland, 1959.

THERMOGRAVIMETRICAL ANALYSIS

CHAPTER 2

61

CHAPTER 2" THERMOGRAVIMETRY

2.1 Introduction

Thermogravimetry (TG), the technique in which the mass of a sample is monitored against time or temperature, is performed with a ThermoGravimetric Analyser (TGA) or thermobalance. Recently, a survey of this technique and the available commercial equipment was given by Wunderlich [i]. Important differences between the balances are- - the type of balance, vertical or horizontal furnace systems

(the horizontal furnace TGA needs a correction for the influence of the thermal expansion on the length of the balance-arm),

- the sensitivity in combination with the maximum sample weight (a typical example is the Perkin Elmer TGA-7 with a sensitivity of 0.0001 mg and a maximum sample weight of 200 milligramme.

- the temperature range and the temperature accuracy. The sample mass determination and the sample temperature measurement of the TGA has to be calibrated using calibrated weights and the ferromagnetic transition (Curie) temperatures of calibration metals.

A Perkin Elmer TGA-7 (vertical furnace) system is used for the TGA experiments described in this chapter. The balance is purged with 40 ml/minute of nitrogen. A second nitrogen purge gas stream of 20 ml/minute is entering the system via the sample purge entrance, see Figure 6.5. Both gas streams purge the furnace part of the TGA. The sample purge can be switched from nitrogen to air with the aid of a software controlled gas-selector.

The automated two-point temperature calibration procedure is performed with the calibration standards alumel and perkalloy. Subsequently, the Curie temperatures of four calibration standards are measured to check the whole TGA temperature range of interest. A typical calibration result mentioned below shows that a transition temperature accuracy of about • I~ can be obtained using the standard temperature calibration procedure"

alumel - 163oc, Curie temperature = 163~ nickel : 356~ " " = 354~ perkailoy: 597oc, " " - 597~ iron : 787~ " " - 787~

The TG technique is (just like the DSC) very popular in polymer reseach, in particular to study the thermal stability of polymeric systems under application conditions. An example of the straightforward use of this technique for polypropylene (PP) is given in chapter 2.2. A series of TG experiments on PP catalyst systems under special conditions (an oxygen- and moisture-free sample loading procedure) is described in

62

chapter 2.3. A number of TGA applications in combination with FTIR and MS analysis of the released purge gas is, finally, mentioned in chapter 6.

2.20ligomers content and thermal stability of polypropylene

2.2,1 The non-isothermal thermal stability determination A quick impression of the thermal stability of a polymer is obtained by measuring its mass with a TGA as a function of the temperature at a constant heating rate. Figure 2.1 shows the results of four TGA scans on ~'P samples using a heating rate of l~ These results illustrate that" - the thermal stability of PP considerably decreases if oxygen

is present and, - that the added stabilisation system only works under inert

conditions. These non-isothermal TG experiments offer a simple and relative quick analytical possibility to compare the thermal stability of different polymers or different polymer batches.

A polymer is thermally stable untill the decomposition process starts. Two (main) types of thermal decomposition processes are usually recognised for polymers, chain depolymerisation and random decomposition. Chain depolymerisation is the release of monomer units from a chain end or at a weak link and is essentially the reverse process of polymerisation. It is often called depropagation or unzipping. Random degradation occurs by chain rupture at random points along the chain, giving a disperse mixture of fragments. These two different processes may occur separately or in combination; the latter case is rather normal [2]. Both processes cause sample mass losses which can be measured with a TGA.

The thermal stability of a polymer is often expressed by its Td(0.5)-value. This is the temperature at which the loss of mass during pyrolysis (at a constant heating rate) reaches 50 % of its final value. This Td(0.5)-value is measured in an inert atmosphere and is, according to Van Krevelen [2], determined by the polymer's chemical structure. It is important to realise, however, that the physical properties of a polymer are changed/decreased significantly at the moment that the Td(0.5) temperature is reached. We prefer, therefore, to characterise the thermal stability of a polymer by the Td(o)-value i.e. the temperature at which a loss of mass during heating just starts. The determination of the Td(o)- value can be hampered, however, by mass loss effects due to a low molecular weight fraction (oligomers) and/or residual monomer/solvent in the polymer sample. Commercial PP grades, for example, always contain a small (0.I %wt. - 0.3 %wt.) oligomer fraction (C6 - C39).

A series of isothermal TG experiments was used to determine the oligomer content of a commercial PP grade. In addition, the Td(o) -value of pure i.e. non-stabilised PP was determined.

e,i~-

I-- ./"

,/ / !

/ I

/ I

63

~1 ~ W n-

N m

-i "r r~_

Q.W

m 0 Z

r 7

W

0 Z w

Lu "r

N

0 I-- _

i |

--

,.,

= ......

- .

..

..

..

..

.

IHrdl3M

%

U

o o ILl

,-- F-

o

Z o~

""

.....

I

_-J Q~ c

56

i

w

i; i w

R R B

i

64

2.2,2 The isQ~hermal thermal stability determination. Isothermal mass/time curves of non-stabilised PP powder samples were measured during I000 minutes at temperatures between 160~ and 280~ The PP TGA samples (about I0 mg.) were flushed with nitrogen during one hour at 30~ before the experiment was started.

Oligomers and polymer are separated, during these TGA experiments, on basis of the boiling temperature (and thus the molecular weight) of the different oligomer fractions. The total amount of oligomers determined during these isothermal TGA experiments will increase, therefore, with increasing measuring temperatures.

The mass/time curve measured at T(isothermal) = 160~ shows a non-linear mass loss effect during the first 250 minutes, see Figure 2.2. No further mass losses were detected during the remaining 750 minutes of this experiment. This mass loss effect of about 0.3 %wt. is assumed to be caused by the evaporation of an oligomeric fraction.

The TGA experiments at T(isothermal) ~ 190~ show a continuous, nearly linear with the time, decreasing sample mass after the first (non-linear) mass losses due to evaporation of the oligomers fraction. The slopes of the linear part of these curves increase with isothermal measuring temperatures. This effect is thought to be caused by the thermal degradation of the polymer matrix. The mass/time curves were extrapolated, subsequently, as indicated in Figure 2.2 to determine the oligomers fractions [wl, w2, w(n)]. The curves in Figure 2.2 clearly illustrate that the extent of the oligomers fraction increases with increasing measuring temperatures. The oligomer content values and slopes of the curves which are thought to represent the rate of the' thermal degradation process of the PP matrix are listed in Table 2.1.

Nib samples (i.e. the extruded and stabilised material) were investigated, subsequently, in the same way. The results of this series of experiments is listed in Table 2.2. The oligomer contents measured for both the powder and the nib samples are plotted as a function of the isothermal measuring temperature in Figure 2.3. The oligomer content of the non- stabilised powder sample increases nearly linearly with the temperature. The oligomer contents measured on the (stabilised) nib samples are lower and scatter considerably more than the powder sample values. These values have to be lower due to evaporation losses during the nib extrusion process. Differences in the cooling rates of single nibs after the extrusion process might be the reason for the considerable amount of scatter. The broken line indicates thus a kind of maximum oligomer content for nib samples.

The loss rates due to thermal degradation of the PP matrix are plotted in Figure 2.4 as a function of the reciprocal absolute temperature. The first detectable loss rate value was measured for the non-stabilised PP sample at T(isothermal) = 191~

I00.099. 9 I ................

g9. 8

. . . . . l ~z

-100.0

- 99.9

T ( l s o t , h . ) - t68~ . . . . . . . . . . . . . . . . . . . . . . . . - - . . . . . . - - _ _ _ _ _ L

99. 8

99. 7

o~" 99,

" " 99,5 r -

---~ U

~: 99,4

9g. 3 -

99. ~4

9 9 . ] -

99.0 -

r

98, 9 - I 0 .0

t 3

T ( t s o t h . ) - 198 ~

T ( t = o t h . ) - 228~

, %

. . . . . I I ...... 1 I I i ' 1 ' 1 1 100. 0 200. 0 300. 0 400. 0 500. 0 600. 0 700. 0 800. 0 900. 0

Time (minutes)

99. 6

99. 5

99. 4

99. 3

- 99.2

- 9 9 . 1

" 99.0

- 98.9 1000. 0

o~

v

r -

Figure 2.2 Isothermal TGA mass/temperature curves of a PP powder sample (non-stabilised)

66

Table 2.1 Results of isothermal TGA experiments on non-stabilised PP powder samples

~,..~[_ ,, . . . . .

T ( isoth. ) oligomer thermal degradation

211

221

oC fraction w (n) , %wt.

,

0.288

loss rate, % wt./s

161 0.0 . . . . .

171 0.353 0.0 = , ,

181 0.278 0.0

191 0.413 6.9E-7

201 0.467 1.4E-6

232

251

0.463 1.4E-6

2 .IE-6 0.585

0.675

0.783

4.9E-6

1.0E-5

Table 2.2

T ( isoth. )

oC

171 ...... , ,

182

192 , ,

202

222

242 , =

250

262

270 ,

277

Results of isothermal TGA experiments on stabilised PP nib samples

. . . . 1 1 I J I I - - 7 i~ I

m

oligomers fraction w(n), % wt.

0.282 , ,

0.286

0.350

0.261 , , ,

0.423 .

0.534

0.425

0.568

0.496

0.412

thermal degradation lOSS r a t e , % wt./s

0 . 0

0 . 0

0.0

7.4E-7 , ,, ....

9.0E-7

7.8E-7

1.5E-6 ,, ,,

4.7E-6

1.6E-5

3.3E-5 �9 7

67

0.80

Figure 2.3 Oligomer content/temperature relation (nonstabilised powder/stabilised nibs)

+ n o n - ~ s tab. s tab. mibs

-I-

0.70 +

c 0 o

E 0

0

0.60

0.50

0.40

0.30

-t-

J _r

J

A

-I-

J J

A

J

A

J J

A

J

A

A

0.20 I 160

_ _ I

180 I

2OO I

2 2 0 2 4 0 I

26O 28O

T(isothermal), ~

cl)

(1)

I I

0 0

-q

O"l

T '"

i I

'l I'

0 CO

-

_,.l

,

LO

- 0

o LO

-

0 o ::::I

" IX.

) b-

0 IX)

...t

. 0 ...l

. "

C)'l

Ix,)

0

0 b 0 lo

ss r

ate,

%w

t./s

0 0

! "

i t

! lr

i !

i i

j I::>

"

+

i>

..

..

.

�9

! II

--I'T

1

(Ii

(::

+ -~

"" (1)

(i}

!~

""II 3~

o)

::3

~-

~ (tO

_,~,..

0 ...,

..

0 (I}

0 [>

.-~

"1

3

(~

c o

.,.,

...

0 0 0 0

O'i (30

69

Hence, the Td(o)-value of pure (non-stabilised) PP is about 190~ It is also clear that the small amounts of oligomers presen t in these samples do not hamper a proper Td(o)-value determination. This Td(o)-value of 190~ increases to about 240~ due to the addition of the stabiliser system (ionol and irganox). The activity of this stabiliser system is not only temperature but also in time limited. Figure 2.5 shows that this stabiliser system stays active (at 250~ for about I000 minutes i.e. long enough to withstand the thermal treatments during all different PP processing procedures.

Usually, non-isothermal TGA experiments are used to determine the 'standard' Td(0.5)- and Td(o)-values. The non-stabilised PP sample was measured, therefore, at different heating rates in order to match a non-isothermally determined Td(o)-value with the isothermally determined Td(o)-value of 190~ This match was obtained, see Figure 2.6, for a heating rate of 0.1~ The slope of this mass/temperature curve increases due to the evaporation of oligomers between 75~ and 184~ and clearly decreases then in the temperature region between 184oC and 204~ The presence of the latter temperature region indicates that there is sufficient separation between the oligomer evaporation process and the start of the polymer matrix degradation. The Td(o)-value of PP determined in this way is estimated at 194~ and is called the semi-static Td(o)-value.

2.3 The TG analysis of a PP catalyst system

2,3.1 A 'plastic wrapped' TGA TIC14, the active component of a Ziegler-Natta PP catalyst system, has to be coordinated on the MgCI2 carrier material together with an internal electron donor like di-isobutyl phthalate (DIBP). The close presence of TiCl4 and DIBP might also result, however, in the formation of a TiCI4.DIBP complex, which would influence directly the catalyst activity.

Different analytical techniques like NMR and IR have been used to study these catalyst systems, but they did not resolve the doubts about the state of the TIC14, the DIBP and possibly the TiCI4.DIBP complex in these catalysts.

Terano et.al, used TG experiments to investigate a TiCl4/ethyl benzoate system (EB) [3]. Their results suggested that the TiCl4 and EB exist seperately on the MgCI2 surface. Based on these results, a series of non-isothermal TGA measurements was performed to study the TiCI4/DIBP system in the same way. A 5~ heating rate was used (Terano used 17oC/minute) to improve the resolving power of this technique. Besides, the TGA sample loading procedure had to be optimised.

These types of catalysts have to be handled in an moisture and oxygen free atmosphere and are stored in a dry-box. Special precautionary measures are, therefore, necessary for the TGA

100.0 -

gg. 5 -

gg. 0 -

98 .5 -

98 .0 -

97 .5 -

97 .0 -

96 .5 -

e..

. = . .

PP NIB S A M P L E s t a b i l i s e d

PP P O W D E R S A M P L E n o n . s t a b i l i s e d

T ime (minutes)

0 .0 500. 0 1000. 0 1500. 0 2000. 0 2500. 0 3000. 0 3500. 0 4000. 0

Figure 2.5 The mass/time curves of a non-stabilised and a stabilised PP sample during isothermal TGA experiments at 250~ (nitrogen atm.)

lO0. 0

99. 5

99. 0

98. 5

g8. 0

gT. 5

97.0

g6. 5

P E R K I N - E L M E R 7 Ser ies T h e r m a l Ana l ys i s Sys tem

100.0 - ' . . . . .

99.8 -

9g. 6 -

9g. 4 -

o~ 9 9 . 2 -

9g.o-

.,m, QI

N 9 a . a -

98.6 -

98.4 -

98. 2 -

98.0 -

50.0

i,i i , , , , , , , , , , / , , , , , , , , , ,,,,,, , , , , , , , ' , , , , , | i

TCA F i l e Nome: t2472 " Somple We19ht: 10.714 m 9 Sun Hot 31 1 4 . 4 9 : 2 3 l g g l ~ ' ~

U n s t o b ~ l I s e d PP powder" ex P e r n l s CPO

nitrogen atmosphere TEMP 1: 30.0~ TIME 1: 0 .0min RATE 1: 0.1~ TEMP 2: 450.0~

184[ ~ 2@41 ~ / , , . . . . . t . . . . ~ I .... ~ ' ' - I . . . . . . . . . . . . I ' I ' "

75. o ~oo. o ~25. o ~5o. o ~75. o 2oo. o 225 .o 250. o 27~.o 3oo. o

Figure 2.6 PP TGA experiment. Heating rate 0.1~

Temperature (~

, , , 3

72

sample loading operation.

The TGA sample pan was filled with about I0 mg. of catalyst sample and placed in a closed, nitrogen flushed, small bottle. This happened in the dry-box used to store the catalyst. The TGA sample pan was transported in this bottle to the TGA sample loading table. The whole TGA balance system was packed, subsequently, in a big polyethylene bag (Sigma A3533 Atmosbag 36" x 51" x 58", provided with two gloves) This bag was closed as good as possible around the electrical cables and blown up with an excess of nitrogen. The relative humidity in the bag decreased in about one hour from 50 - 70 % to less than 5 % due to this nitrogen purge. The catalyst sample was then, using the gloves of the Atmosbag, taken out of the bottle and placed in the TGA hang-up wire. Subsequently, the TGA furnace was closed and the measurement was started.

2.3.2 TG analysis of a MaCl2-suppor~ed. TiCI4/DIBP catalyst _ v

Figure 2.7 shows the TG mass/temperature curves of the reference systems MgCI2, DIBP and a mixture of MgCI2/DIBP. The decomposition of MgCI2 starts at temperatures > 400~ The total mass loss between 420~ and 730~ was 17.7 %wt. The DIBP as such evaporated completely between II0~ and 250~ The MgCI2/DIBP mixture mass/temperature curve clearly shows a two- step mass loss effect due to evaporation of the DIBP as such and due to evaporation or possibly decomposition of DIBP absorped on the MgCI2 surface. A mass loss effect of 25 %wt. is measured between II0~ and 300~ (evaporation) while 11.3 %wt. weight loss is measured between 510~ and 584~ (evaporation or decomposition). The total weight loss due to DIBP release (after correction for the MgCI2 weight losses) is thus 36.3 %wt. This value fairly agrees with the DIBP content of 34.7 %wt. reported for this system.

Two other reference systems, MgCl2-supported TIC14 and TiCI4.DIBP complex were measured in the same way. The TIC14 decomposes in two steps with DTGA minima at I19~ and 221~ The total weight loss at 420~ (the starting temperature of the MgCI2 mass loss process) was 18.5 %wt. The calculated weight loss due to a conversion of the TIC14 to Ti is 20.0 %wt.

The main mass loss effects of the TiCI4.DIBP complex occur between 128~ and 288~ with at least two DTGA minima at about 184~ and about 220~ The main evaporation process of DIBP both from TiCI4.DIBP complex and from the MgCI2/DIBP mixture thus occurs in the same temperature region which was disappointing. On the other hand, the TiCI4.DIBP complex shows a small but clearly present effect at about 400~ which is not present in the MgCI2/DIBP mixture results. Besides, the MgCl2/ DIBP mixture shows a strong effect between 510~ and 584~ (Figure 2.7) which is hardly present in the TiCI4.DIBP complex results.

~'C',-"c"---~ ~ ' : ':L=-i - ' ~"-----~ 100. 0 ~ : L " 1 e ""- ' -~-~

go. 0

8 5 . 0

~ -I HgCl 2 / D I B P m i x t u r e ..; BO. 0

~ _ 1 [ ,

= 75. o 3 " ~8 c

70. 0 -

6 5 . 0 -

60. 0 -

55. 0 -

50. 0

7 5 . 8 ~.

D IBP, IB8 ~'.=t. 1o==

betueen t 1( ] -25( ] C.

Commercial M g C 1 2 ~

1?.7 ~=t , , loss between 42B-730 C

se4 c 5 5 . 9 ~

,1~ . . . . . . . . . . . . . . . . . . l j[ l . . . . . 1 . . . . . . . . . 1 ' ' l I . . . . I I - -

1oo. o 200. o 300. o 400. o 5oo. o son. o ~oo. o Figure 2.7 Temperature (~ Non-isothermal TGA mass/temperature curves of DIBP, MgC12 and a mixture of MgC12/DIBP measured in a nitrogen atmosphere

74

Figure 2.8 shows the TG mass/temperature curves and the corresponding DTGA curves of the MgCl2-supported, TiCI4/DIBP catalyst in duplo. This catalyst contained 15.8 %wt. of TIC14 and 20.8 %wt. of DIBP. These two mass/temperature curves have roughly the same shape as the MgCI2/DIBP mixture mass/ temperature curve shown in Figure 2.7. This is an indication that an important part of the DIBP in this catalyst sample is present as DIBP and not as TiCI4.DIBP complex. The "blown-up" DTGA curves in Figure 2.9 illustrate that the TiCI4.DIBP complex DTGA minimum at 390~ is not present in the catalyst samples. These results are, therefore, suggesting strongly that the DIBP is present as DIBP in this catalyst system and not as TiCI4.DIBP complex.

References

i. B. Wunderlich. Thermal Analysis, Academic Press Inc., New York, 1990.

2. D.W. van Krevelen- Properties of Polymers, Third edition, Elsevier, Amsterdam, 1990.

3. M. Terano and T. Kataoka, Makromol. Chemie, i__~, (1987), p. 1477 - 1487.

100. 0

95, 0

go. 0

85. 0

A

8o.o

- - 7 5 . 0 - J:: . , , . .

(9

7o.o-

65, 0 -

60. 0

55, 0 -

50, 0

k

1 5 2 8 ~

/ I # //

I ! / I

234~ . . . . . . . . ~ . . . . . . . . . . . . I . . . . . . . . . I ..... I i . . . . . . i . . . . . . . t

I00 . 0 200 .0 300. 0 400. 0 500. 0 600, 0 700, 0

Figure 2.8 Temperature (~ Results of two non-isothermal TGA measurements on PP catalyst BR900718 in a nitrogen atmosphere

�9 -0, I

" -0 . 2

" -0 . 3

" -0 . 4

" -0 . 5

- 0 . 6

-0~ 7

- 0 . 8

- 0 . 9

- 1 . 0

.,,1 U1

- 0 . 1

-0. 2

-0, 3

.c_ -0. 4 E o~ "-~ -0. 5 ID

~ - 0 . 6 .~_ L _

E3 -0 .7

- 0 . 8

-0. 9

- 1 . 0

-1 .1

\ |

\

t

t

I

i/ / I /

r i !

f i i !

t

|

.J ' / \., [

390~

I

;

I

15.BY. T ICI4 28. B~. DIBP

I

4.BX T tCI4 �9 22.5~ DIBP

5aeoc

-0. 1

-0. 2

-0. 3

-0. 5

-0, 6

-0. 7

-0. B

-0. g

.0

-1 .1

,..3

~o0. o 2oo. o a~o. o ,od. o sod. o coo. o 7od. o �9 Temperature (~

Figure 2.9 DTGA minima of two PP catalyst samples and a TiC14.DIBP complex sample measured at 5~ in a nitrogen atmosphere

THERMODILATOMETRY

CHAPTER 3

77

Chapter 3- THERMODILATOMETRY

3.1 Length dilatometry (TMA)

3,1.~ Introduction Thermodilatometry, the technique in which the dimensions of a sample are monitored against time or temperature, is performed as length dilatometry or as volume dilatometry. Length dilatometry uses a thermomechanical analyser ~TMA) which measures the sample length as a function of the temperature or time while the sample is held under a small, constant compression force. Volume dilatometry, performed with the classical dilatometer, uses a liquid (for example mercury) to measure the volume change of a sample as a function of the t~mper~ta~eor time. Due to its simple, easy to operate, technique is length dilatometry far more popular in polymer research applications than the complicated, time-consumin~ volume dilatometry. TMA equipment is, nowadays, available from several manufacturers. Recently, also dynamic load thermomechanical analysers (DLTMA) became commercially available. A survey of the TMA/DLTMA techniques was given by Wunderlich [i].

The TMA technique can be used for Tg-value determinations, resin cure studies, penetration experiments or orientation effect determinations. The most important application is thought to be the linear thermal expansion coefficient (l.e.c.) determination of engineering polymers. An example of this application is given in chapter 3.1.2. The results of a polymer shrinkage experiment monitored by TMA are described in chapter 3.1.3.

The TMA used for the l.e.c, determinations described in this chapter is a Perkin Elmer TMA-7. This TMA is purged with 60 ml/minute of nitrogen. The system is operated at a heating rate of 2~ while a small (I0 mN) compression force is put on the sample. The length measurement is calibrated using (calibration) standards; the automated two-point temperature calibration program is performed using indium (156.6~ a S1215 rubber sample [Tg(midpoint)-value = -36oC].

3.1.2 The linear thermal expansion coefficien~ determin@tion of filled polyketone systems

Engineering polymers are often filled with glass fibres or other filler types to improve certain mechanical properties, like stiffness and thermal expansion. The thermal expansion of polyketone* samples filled with different filler materials

Aliphatic polyketone based on carbon monoxide, ethylene and a small amount of propylene, commercialised by Shell under the trademark CARILON Polymer (PK-EP), see Chapter 9.

A

E E c: 0

~ u) r - m c2. x

LU

2.33

2.32 Sample Height: 2.313 mm

| , h e a t , l n g

2.31

2.30

2.29

2.28

2.27 2 , c o o I I ng

2.26

2.25 3 , h e a t i n g

2.24 1- 22:3 I

-50.0 I I ! I I I I !

-25.0 0.0 25.0 50.0 75.0 tOO 0 t25.0 150.0

T e m p e r a t u r e (~ Figure 3.1 The length change of a polyketone sample during three subsequent TMA scans

79

was measured to study the influence of the shape of the fillers on their l.e.c, depressing efficiency.

The TMA samples of 5 x 5 x 2 mm are machined from compression moulded (i00 x i00 x 2 mm) or injection moulded (60 x 60 x 2 mm) samples. Stresses, 'frozen-in' during the injection or compression moulding procedures, hamper a straightforward l.e.c, determination. Figure 3.1 shows the length/temperature relation of an injection moulded polyketone sample measured in the z-direction i.e. the plane perpendicular to the sprue xy- plane. The sample shrinks about 3 % in this direction due to the release of these 'frozen-in' stresses during heating and cooling between -50~ and 150~ The length/temperature relation measured during the third heating scan is nearly equal to the cooling scan result. This indicates that the 'frozen-in' stresses are completely released. It will be clear that a proper comparison of the effect of different filler systems on the thermal expansion is only possible on basis of results measured on stress-free samples.

Addition of fibre-shaped fillers introduces anisotropy in the material. Characterisation of the thermal expansion properties requires therefore measurements in all three directions. Measurements in the x-, y- and z-direction of a pure, non- filled material resulted in the following average l.e.c. values at 20~

Polyketone �9 1.lIE-4 • 0.02E-4 K^-I, x-direction l.e.c, at 20~ �9 1.09E-4 • 0.02E-4 K^-I, y-direction

�9 1.09E-4 • 0.02E-4 K^-I, z-direction (The sprue-plane is called the xy-plane. Results measured on four different TMA samples taken from an injection moulded sample, crystallinity 44 %wt.)

The polymer as such is considered to be isotropic i.e. the l.e.c, values in the three directions are equal, 1.10E-4 K^-I. Polyketone is, however, a semi-crystalline polymer and the l.e.c, will be influenced by the extent of the crystalline phase :

Polyketone : 1.02E-4 K^-I, crystallinity 54 %wt. l.e.c, at 20~ �9 1.10E-4 K^-I, crystallinity 44 %wt.

- 1.16E-4 K^-I, crystallinity 36 %wt. (average l.e.c, values in the x-, y- and z-directions)

The l.e.c, values of a series of eight polyketone samples containing 30 %wt~ of six different filler materials were measured. The l.e.c, values measured in the x- and y- directions proved to be nearly equal pointing at a random distribution of the fibres in the xy-plane, the average values are listed in Table 3.1. A grafical representation of these data is given in Figure 3.2.

80

(/)

L Q)

C"

\ L

"~"

(.- r

L ~

~ ~

~

-0--

0

0 O--

~ L

~_~

0 .~-

r U

3 E

tfl

+ x

4 I

~

l-X---I / / /

/ *--I---4

i

I---X--I

- ~K

=

.... J

: ~'

I

I-- ..I- --.I

t .

1 / _

. ,

I ,

I i

I ~

I ...

, I

, __

I ,

.. t

03 CO

r,,.

f.D

IS') q-

or) ou /

L_

0 i I .

D

4k~

r 0 L- 0

,u

-J

I,I

F-

0.0

0 ~-

&

Nk--

,iii,

r m

If,,

~'~E

I t

�9

_

0o0~ le uo!loeJ!p ~ u!

(NIL 'g-3x) "geoo uo!suedxe .m

eu!l

0I I- 0!

E .e~

r- ,-.=

== .Q

.i..

LL

0 iI

Ol

I,- 0!

9-

81

Table 3.1 Linear expansion coefficient values at 20~ of polyketone samples with different filler materials

filler type

30 %wt. . . . . . . .

CaC03

kaolin

wollastonite

mica

short glass fibres

long glass fibres

average longest

dimension, micron

........

1.5

2

4O

20 80

125 150

7000

x-/y-direction average l.e.c.

at 20~ K^_I

. . . . .

8.7E-5

8.2E-5

8.1E-5

7.5E-5 6.0E-5

..

4.7E-5 5.7E-5

z -direction l.e.c, at

20oc, K^_I

.,

9.4E-5

I.IE-4

I.IE-4

1.2E-4 1.5E-4

1.6E-4 1.6E-4

3.3E-5

L1

1.9E-4

Quite different filler materials with average longest dimension values varying over four decades fairly fit the drawn straight line. This indicates that the average fibre length determines (at a constant filler concentration) the l.e.c, in the xy-plane. Scatter in the data is expected to be proportional with the filler dimension inhomogeneity (wollastonite) or with the sensitivity to the processing procedure (short glass fibres during injection moulding), see Figure 3.2.

The linear expansion coefficient values shown in Table 3.1 and Figure 3.2 illustrate that filler addition results in composite systems with clearly decreased linear expansion coefficient values in the xy-plane. It also introduces, however, anisotropy in such a composite system i.e. an equal or even higher linear expansion coefficient value is measured in the z-direction.

3.1.3 Shrinkaue of polvketone and nvlon 6.6 due to moisture

The dimensions of a polymeric component are temperature, time and moisture content dependent. The extent of these dimensional instability effects, especially those related to temperature and moisture content is important if the physical properties of different engineering polymers are compared. The expansion due to moisture absorption of polyketone and that of a alternative system Minlon 13TI (Nylon 6.6 with 33 %wt. mineral filler) were measured in order to obtain such a comparison.

A

E E

v

C 0 ~ u~ r

Q. x

LU

2. i 74 -

2. 173 -

2. 172 -

2 . i 7 1 -

2. 170 -

2. 169 -

2 . i 6 8 -

2. t67 -

2. 166 -

2 . i 6 5 -

2 . i 6 4 -

2. i 6 3 -

2. 162 -

2 . 1 6 ! -

2. 160 -

X!

X2

YI

Y2 ! AY

O. 000 min ,,

9795. i 50 min

2. !73 mm

2.161 am

- 0 . 0 1 2 mm

" ' " 1 I I ' 1 i I ! I I 0 2000 4000 6000 8000

T i m e (m inu tes )

Figure 3.3 The shrinkage of a polyketone sample due to loss by evaporation of absorbed water

! 0000

83

It proved practically impossible to keep the TMA sample immersed in water in the (small) TMA measuring cell during more than twenty days. However, measuring the shrinkage during drying instead of the expansion during water absorption is very well possible with the TMA. This assumes that the expansion (due to moisture absorption) and subsequent shrinkage (due to drying) are completely reversible effects. Witchey et al. showed that this indeed holds for poly[l- (trimethylsilyl)-l-propyne] in contact with n-nonane [2].

Both TMA samples were first stored at 22~ in distilled water until an equilibrium water saturation was reached. The polyketone reached its equilibrium water saturation of 2.35 %wt. in about twenty days. The Minlon 13TI sample needed forty days to reach its equilibrium water saturation of 6.12 %wt. (based on to~al sample weight). The 'wet' samples were put in the (N2 purged) TMA and the length decrease due to moisture loss was measured as a function of time. The average temperature during these experiments was 22~ • 2~ (this 2~ temperature scatter was mainly caused by a day/night temperature difference). After the TMA experiments both samples were completely dried by storage at 50~ in vacuum, to determine the (residual) moisture concentration.

Figure 3.3 shows the shrinkage effect of the polyketone sample due to a decrease of the moisture content from 2.35 %wt. to 0.08 %wt. in about seven days (the day/night temperature differences are clearly 'modulating' the experimental results). The smoothed results of the experiments on both samples are shown in Figure 3.4 and can (assuming a linear relation between moisture content and amount of shrinkage) be expressed as:

polyketone �9 0.25 % shrinkage/percent of moisture loss, Minlon 13TI- 0.52 % shrinkage/percent of moisture loss.

SRMPLE THICKNESS DECRERSE, 2.78

48Ttu=t,ercont,esnt,z G. 12 ~ut,.

2

2 tO

l .80

I . 5 0

I .28

.88

.G8

.38

8.88

+

X +

+ \ \ +

Figure 3.4 Shrinkage of completely water saturated polymer samples as a function of the drying time (TMA experiment, N 2 atmosphere, 22~

( , ~ t t e r con t , ont. z 2 .35 ~wt,.

X,, .X X

" X p o l y k e t o n e x ~

i J i ~) oj

6.6

~ ~ +

~ X ~ ~ wmtercont ,ent , : 8 . 0 8 ~ u t ,

co c)o (s) oJ .. , . 4 . , " 4

_ I

w a t , e r + cont,ent, :

i _ i J i . , i ' i

o

( t Cs ] ) ^ ( 1 / 2 )

85

3.2 Volume dilatometry

3.2.1 .... Introductio~ Polyisoprenes (IR) polymerised with a Ziegler catalyst have cis contents of about 98%. The last few per cents of non-cis polymer, which makes IR different from natural rubber (NR), are very important for the physical properties of the vulcanizates. The uncertainty in the NMR analysis, amounting up to • 1%wt. was the reason to investigate the possibilities of dilatometric crystallisation rate measurements.

Mitchell [3] demonstrated the use of a recording dilatometer for the determination of the crystallisation rate. He used as the rate characteristic the time necessary to reach 50 % of the ultimate change in volume due to crystallisation. This time is called the 'crystallisation half-time value'.

The potential of the crystallisation half-time values to distinguish between IR batches from various sources and made under different conditions, was studied. A continuously recording dilatometer was built for these experiments. This apparatus, the measuring method and some results are described in the following sections.

3,2,2 The volume dilatometer Figure 3.5 gives a schematic diagram of the volume dilatometer . The cell containing the rubber sample is made of stainless steel and consists of two parts which are screwed together (Figure 3.6). The lid, which contains a deep thermocouple well, is provided with a metal-to-metal seal; as an extra safeguard against leakage an O-ring is compressed within an annular chamber outside the cell. In order to prevent blocking by the rubber of the connection between cell and capillary, the latter is attached to the bottom of the cell.

Filling the system with mercury (the measuring liquid) and measuring is done via a system of glass capillaries and stopcocks. This permits to work at different sensitivities simply by opening one of the stopcocks, thus giving the mercury access to a capillary which acts as a shunt with respect to the measuring capillary.

The measuring capillary tube is metallised at the outside. This metallic layer acts as the fixed electrode of a variable capacitor, the variable electrode being the mercury inside the glass capillary tube. A wide tube which is metallised except for a double slit is electrically connected to the mercury and acts as shield. Details are shown in Figure 3.7.

The variations in capacity are detected with a highly sensitive capacitance bridge assembly (Wayne-Kerr, Autobalance Universal Bridge B641), allowing the zero capacity of the system to be compensated, so that the variation in capacity is measured with maximum sensitivity. The bridge offers an output voltage proportional to the capacitance variation which is

86

MEASURING CAPILLAR~Y

GUARD

HIGH POTENTIAL ELECTRODE

BALL JOINTS

ACCESS TO VACUUM PUMP

t

\ \ ,< ~ \

X,

FILLING , CAPI LLAR_Y.

THREE-WAY STOPCOCKS

AL LOW POTE, NTI ELECTR, 0DE

COOLING WATER I N ~

COOLING WATER OUT

GASEOUS N 2 I IN . ~--,

i

DRAIN , .

HEAT SINK PERSPEX HOUSING

DRY NITROGEN A T M O S P H E R E

',S,~ j I

V [____ ' ,

I STAINLESS STEEL MEASURING (~ELL

I ALUMINIUM,THER- MOSTAT BLOCK

COOLING (PELTIER) i ELEMENTS

ISOLATING SUPPORT

Figure 3.5 Schematic arrangement of apparatus for volumetric crystallization measurements

87

Figure 3.6: Sample holder

88

Figure 3.7: Measuring capillary

89

recorded as a function of time together with the sample temperature. The drift of the Wayne-Kerr Bridge proved not detectable by the recorder over a period of more than one week. The detection system is trimmed to ensure that the volume change during crystallisation corresponds to about full-scale on the recorder. This results in a sensitivity in (delta V)/V of the order of 10E-4.

The temperature is controlled electronically with a Cu/Const thermocouple as the sensor and Peltier elements as power drains, which surround an aluminium block containing the steel cell. The measuring cell is sensitive to outside temperature variations via its metallic connection to the measuring capillary. However, this sensitivity largely eliminated by the fairly massive heat sink attached to the metal capillary inside the perspex housing, whose interior temperature varies little due to the circulation of cooling water on the 'hot' side of the Peltier elements. In general, a temperature stability of better than 0.2~ was measured. The current through the Peltier elements is reversed during the heat pretreatment of the samples (see 3.2.3). The Peltier elements act in that case as heating elements instead of cooling units. Using the Peltier elements in this way, the temperature region available ranged from -40~ up to 85~ Protecting devices are built in to prevent overheating (> 85~ and damage due to interruption of the cooling water supply.

3.2.3 The measurina procedure _ _

The rubber sample to be investigated should not contain any entrapped or dissolved air or other gas. So, prior to the admission of the confining mercury careful evacuation of the cell containing the sample is required. Other precautions to guarantee reliable measuring results are to melt any residual crystallites and to relieve any internal stress in the sample. Mitchell used a procedure of 40 minutes/100~ [3]. Martin and Mandelkern employed a thermal pretreatment of one hour at 60~ [4]. A series of scouting experiments with NR showed that one hour at 80~ was sufficient to investigate NR/IR systems.

The system was thus heated in about I0 minutes to 80~ and kept at that temperature for one hour. Subsequently, the system was cooled in about 50 minutes to -26~ the crystallisation temperature, and kept constant at that temperature. All volume/time curves obtained in this regime initially showed a very slowly increasing slope up to a maximum at a point of inflection and a very slowly decreasing slope in the latter phase of crystallisation, so that the curve seemed to approach an end level asymptotically. The measurements were continued for at least 3 x t(0.5); further prolongation would alter the calculated half-time value only slightly. A series of seven measurements on a NR sample gives an impression of the reproducibility of these experiments-

90

tl/2, h

5 6 -

4 8 -

4 0 -

32 -

24

16

8

&

& NATSYN 400 WITH 2% STEARIC ACIO ~ _ o NATSYN 2200

o I , , ,I i ,,, ! o - e - ~s - 2 4 - 3 z

Crysta l l izat ion tempera tu re , ~ F igure 3 .8 Crys ta l l i za t ion ha l f - t ime as a func t ion of c rys ta l l i za t ion t e m p e r a t u r e

91

NR, half-time value �9 141 + 3 minutes spec. volume decrease- 0.022 + 0.001 cm3/g

(Crystallisation temperature -25.5~ + 0.2~

3.2.4 Isothermal crvst~llisation of IR rubber systems The crystallisation-half-time value [t(0.5)-value] was first measured at temperatures between -3~ and -29~ to determine the optimum crystallisation temperature. These measurements were performed on three samples- a natural rubber and the Ziegler polyisoprene systems Natsyn 2200 and Natsyn 400 (containing 2 %wt. stearic acid as crystallisation promotor).

The results are plotted in Figure 3.8. A smooth curve can be drawn through the measuring points for natural rubber and the Natsyn 2200. The Natsyn 400 results scatter more than the results of the two other samples, probably due to the stearic acid addition. These curves show that optimum crystallisation conditions are met in the temperature region from -24~ to - 26~ Besides, the t(0.5)-value temperature sensitivity is low in this region.

The results of a series of t(0.5)-value determinations are listed in Table 3.2 along with the spectrometric data available (proton NMR at 220 MHz. and IR). There is no doubt that a variation in half-time values from I0 to 25 hours for samples of 99 %wt. cis and from 11.7 to 54 hours for those of 98 %wt. cis means that the discriminating power of the crystallisation data is far in excess of that of ~R.

Figure 3.9, finally, shows that mixing of NR with the non- crystallising polyisoprene IR 305 (cis content < 90 %) retards the crystallisation process of the NR phase and gives a nearly proportional decrease in the final volume change. Mixing of NR with carbon black is causing the same effects. The timescale of the NR recrystallisation process strongly increases due to blending the NR with IR 305 and carbon black. This sensitivity of the NR recrystallisation for compounding agents offers the possibility to use this technique for mixing efficiency studies of NR based rubber compounds.

References

I. B. Wunderlich- Thermal Analysis, Academic Press Inc., New York, 1990.

2. L.C. Witchey-Lakshmanan et.al., J. of Pol. Sc.- Part B, Polymer Physics, 31, (1993), p. 1545- 1553.

3. J.C. Mitchell, Techn. Rep. No. 278-60, Shell Development Company.

4. G.M. Martin and L. Mandelkern, J. Appl. Phys., 34, (1963), p. 2312.

92

Table 3.2 Results of t (0.5) -value determinations on commercial and experimental polyisoprene rubbers

I - , , I , I I , I I | ' i ,I i J

sample

natural rubber

Z 59

I PIK56, Bayer

Z 58

N a t s y n 2000

,! Amerip____~ol SN600 [ z sG____L____

Natsyn 2 2 0 0

Natsyn 4 1 0

Z 64

, ,, ,, , , , ,, , , r ,

cis 1,4 3,4 content, content, NMR, %wt. IR, %wt.

i I00 0.00

99 0.5O

99 n . d . ! 99 0.60 i

99 0.65

99 i 0 . 4 5

,,, ,,,

t ( o . 5 ) - value,

2 . 5

1 0 . 0

1 8 . 0

20.0

21.0

25.0

11 ._____/7

12.7

19.2

47.5

54.0 , , , , . . . . . , , ,, ,, , , , ~ . . . . . . . . , , , ,~ ,

n . d . , no t d e t e r m i n e d Z numbers, e x p e r i m e n t a l samples

98 n.d.

98 0.90 ,

98 0.60

98 0.80 ' , i

A V , 10 "=, cm=.g "1

2

1.6

1.2

_ . . _ . f - - - - - - - q

RUBBER

A r . I . . 4 ~ - - " 4 " -

RSS ~ WITH IR 305 80/20

i RSS ]3~ WITH CARBON BLACK 1.21 AFTER 30 HOURS

0.4 RSS 111", IR 305 AND CARBON BLACK

O k 0 4 8 12 16 20 24

Figure 3.9 Effect of compounding agents on the crystallization of natural rubber

28 32 CRYSTALLIZATION TIME, h

DYNAMIC MECHANICAL ANALYSIS

CHAPTER 4

94

CHAPTER 4: DYNAMIC MECHANICAL ANALYSIS

4.1 The standard DMA technique

4.1.I Introduction Dynamic Mechanical Analysis (DMA) is a technique in which the elastic and viscous response of a sample under oscillating load, are monitored against temperature, time or frequency. This technique became well known by the impressive amount of information about the structure of polymers obtained with the torsion pendulum apparatus. The torsion pendulum DMA apparatus is a so-called resonant system i.e. the measuring frequency is not constant. The modern DMA systems are nearly always fixed frequency systems operating at frequencies between about 0.01 and I00 Hz. and in a temperature region ranging from about - 150~ to 300~ A survey of the DMA technique and the available commercial equipment was given by Wunderlich [i].

Hooke's law describes the response of a perfect elastic material to an applied stress i.e.

stress (o)/strain (c) - constant 4.1

This proportionality constant is defined as: E, the elastic or Youngs modulus for tensile deformation, G, the shear modulus for shear deformation and B, the bulk modulus for compressional deformation. These three moduli are interrelated:

E = 2. (I + ~).G = 3. (I - 2~).B

where ~ is the Poisson ratio. Newton's law describes the response of an ideal liquid to an applied stress i.e.

4.2

stress (G)/rate of strain (dc/dt) --constant 4.3

This proportionality constant is called the viscosity (~). The response of a polymeric material to an applied stress shows both an elastic and a viscous component i.e. a polymer behaves visco-elastic. DMA equipment measures dynamically the E or G moduli; the polymer samples are assumed to behave linearly visco-elastic i.e. the stress/strain relation is only a function of time. An oscillating (sinusoidal) strain,

(t) = ~0 sin ~t 4.4

during such a DMA experiment results in a sinusoidal stress-

G(t) = ~0 sin (wt + ~) 4.5

with a phase difference ~ due to the visco-elastic character of the polymer sample. The absolute value of, for example, the E modulus is written as.

95

IEI = crolco i.e. ~o = Eo.I'~I 4.6

Substitution of 4.6 in 4.5 and application of the goniometric summing rule gives:

G(t) = E,.(iEi.sin et.cos ~ + IEi.cos ~t.sin ~) 4.7

if, E' = IEi.cos ~ and E'' = IEl.sin ~ 4.7 can be written as :

~(t) - E'.~(t) + (E''/~).d~/dt 4.8

where- E' = the in-phase, elastic component, see 4.1 and E' '/~ = the out-phase, viscous component, see 4.3.

The moduli can thus be written as complex values:

E* = E' + iE'' G* = G' + iG'' and B* = B' + iB'' 4.9

The tangent of the phase difference is given by.

tan ~ = E''/E', G''/G' and B''/B' 4.10

Usually, the E' or G' modulus and the tan ~ are measured during a DMA experiment; E'' or G'' is then also known, according to 4.10.

DMAmeasurements are intensively used to investigate the amorphous phase transitions of polymers. The results of DMA studies were published by authors like Schmieder and Wolf [2], Nielsen and Buchdahl [3] and Heijboer [4]. Neat polymers, but also polymer blends and polymer systems blended with fillers, plasticisers or impact improvers were investigated by DMA. An example of such an application is given for toughened polypropylene in 4.1.2.

Amorphous phase transitions (like the Tg-value for example) can be measured by DMA but also by DSC (chapter 1) and by TMA (chapter 3). The DMA technique, however, offers the highest sensitivity to detect phase transition effects. This is especially useful for the investigation of secondary relaxation effects and for the determination of very weak glass-rubber transition effects. Figure 4.1 shows the results of DMA and DSC measurements on a (crosslinked) epoxy resin system. The (DMA) E' and the E'' curves as a function of temperature both show the system's glass-rubber transition region at about 100oc. Besides, the presence of a secondary relaxation effect with an E'' maximum at about -60~ is shown by the E'' curve. Figure 4.1 also shows the (DSC) specific heat (Cp)/temperature curve of the same sample. The glass- rubber transition is easily detected by the step-wise change in the Cp/T curve at about 100oc. The higher sensitivity of the DMA technique in comparison with that of the DSC technique is illustrated by the clear absence of any relaxation effect in the low temperature part of the Cp/T curve. An example of detection of weak glass-rubber transition effects by DMA is

96

A

EL v

[u OI 0

J

Dynamic mechanical properties of liquid DGEBA resin cured with HI-PA

+ +Log E' & &Log E"

I0.00 19.50 [ +- +-4"+ + + + + + + +_+~+,,+,

9.O0

8.50 E,~/z~" 8.00

8.00 A

~ ' ' ~ + +--4- 7.50 / Polymer LaDS DMTA /

7.50 I- he~ltir~l rate: 2 C/rain / / . ~ ~ y : 10 Hz 1

7.00 7.00 - 1 5 0 - 9 0 - 3 0 3 0 9 0 1 5 0 2 1 0

Temperature. degrees C

0.75

0.63

0.50

"~_ 0.38 o

d. 0 0.25

0.13

Si3er162 heat/temperature relation of liquid DGEBA resin cured with HHPA

F~rk~ Eml~r D$C2 I~atlr~ rst~ 20 C/rain

,J~'m~le wet a h t: ~K24 rr~

4.+--+--+ I I + ̀+

/ I

/+.-"

0 . 0 0 , , ' , ,

- 1 5 0 - 9 0 - 3 0

I I I I I i .

0.75

0.63

0.50

0.38

0.25

0.13

0.00 30 90 150 210

Figure 4.1 T e m p e r a t u r e . degrees C

DMA/DSC results comparison

tu o~ o _J

c~

o

d (J

97

given for rigid polyurethane foam systems in 4.1.3. Performance and results of dynamic mechanical measurements at ultrasonic frequencies, finally, are reported in chapter 4.2.

4.1.2 DMA analysis of DolvDroDvlene/ethvlene-DroDvlene rubber blends Polypropylene (PP) is often blended with ethylene/propylene (EP) rubbers to improve the impact resistance. This so-called toughened PP (TPP) can be a mechanically blended PP/C2C3 rubber system or an in-situ polymerised PP/C2C3 rubber system. A number of rubber parameters (like concentration, particle size, particle size distribution, crystallinity, molecular weight etc.) determine the ultimate effect of the rubber addition on the impact resistance. DMA is one of the analytical techniques often used to investigate blends of polymers with an impact improver. The determination of the relation between the area of the rubber relaxation maximum as measured by DMA and the rubber concentration is usually a first step in such an investigation. The method to determine this area and the results measured on a series of PP/C2C3 rubber blends are reported below.

These results were measured with an automated torsion pendulum apparatus. A rectangular sample strip of 50 x i0 x I mm. acted in combination with a steel suspension wire and a rotating mass as a visco-elastic spring. The measurements were performed while the sample temperature continously increased at a rate of l~ The storage shear (G') modulus and the loss shear (G'') modulus were determined from the free, damped vibrations (frequency about 0.5 Hz.) according to (4):

G' = ((omega^2 - alpha^2).FJ- CST)/FI 4 .II

G'' = (2.alpha.omega.FJ)/FI 4.12

where omega- 6.2832/vibration time, alpha: log. decrement of the damped oscillation, FJ : moment of inertia of rotating mass, CST �9 torsial stiffness of the suspension wire, FI : geometrical stiffness factor of the

rectangular sample

Figure 4.2 shows the DMA results of a PP/talc (85/15) system represented in the standard way i.e. the log G' and log G'' plotted as a function of the temperature. The loss shear modulus curve shows relaxation maxima at about 60~ (the crystalline phase [u] relaxation) and at about 0"C the amorphous phase glass-rubber [K] relaxation). Blending such a PP sample with a C2C3 rubber results in an extra (rubber) relaxation maximum at about -50~

A reproducible base-line drawing procedure and a kind of peak deconvolution procedure is necessary to determine the area of this rubber loss maximum and that of the original PP relaxation maxima. A computerised deconvolution program for the separation of overlapping DMA loss maxima was described by

98 Figure 4.2 Dynamic mechanical properties of PP filled with talc (15%wt.)

+ s to rage A loss rnodu/us modu/us

O4 E

I e + 0 9 v

::3

"0 0 E

c- oo

0 O0

l e + 0 8 - 9 0

w ~

/k

/k

A

A

\

I ..I I I , . ~ - -

- 5 0 - 1 0 30 70 110

A

A-z~,,AJ

I'- 0

g~

9 0 0._ c. c o~

�9 v

1

Z 3 ro

3 e + 0 7

Temperature, deg. C

99

Charlesworth [5]. We still manually "deconvoluted" the loss maxima of the investigated PP samples by applying Heijboer's approach [6] which is based upon the symmetry of the loss maxima around 1/Tmax. when the loss values are plotted on a linear scale as a function of the reciprocal absolute temperature. The DMA results of the PP/talc sample in Figure 4.2 were replotted in this way, see Figure 4.3. The three partly overlapping PP relaxation effects hampered the drawing of a straight base-line in Figure 4.3 as suggested by Fay et al. [7]. A curve fitting model with two fixed points (1/T = 2.40 and 1/T = 5.40) was used to draw a baseline in a standardised way, see Figure 4.3. This curve fitting model was, subsequently, used to draw a baseline in all the experimental PP/C2C3 rubber loss curves. In Figure 4.4 data are replotted with separated effects for both relaxation maxima (based on symmetry around the 1/Tmax. axis) and after subtraction of the background loss. Figure 4.5, finally, shows the result of this procedure for a PP/C2C3 rubber/talc (70/15/15) blend. The loss relaxation areas of these two samples, determined by integration, were:

PP u-rel. PP E-rel. C2C3 rubber rel. (Nm-2 .K-I) (Nm-2 .K-I) (Nm-2 .K-I)

PP/talc (85/15) 6.7E3 12.2E3 PP/C2C3/talc

(70/15/15) 9.7E3 20.6E3 6.1E3

The rubber addition results, as expected, in the presence of a low temperature rubber loss maximum. The DMAmeasurements also show, however, that a part of the added rubber phase influences the intensity of both PP relaxation effects. Such an effect might be important for the toughening efficiency of the used type of rubber.

A series of PP/C2C3 rubber blends was measured, subsequently, to investigate the relation between the loss modulus areas of the rubber relaxation effects and the rubber concentration. The results of these measurements are listed in Table 4.1 and plotted in Figure 4.6.

The Vistalon 404 and 4608 are amorphous C2C3 rubbers, while Vistalon 5600 contains a crystalline fraction of 4 %wt. The Nordel 1500 and 3391 are typical semi-crystalline C2C3 rubbers with crystalline fractions of respectively 15 %wt. and 12 %wt. Each of these rubbers was added to the PP in three different concentrations i.e. 10, 15 and 20 %wt. 0nly the amorphous fraction of these C2C3 rubbers is assumed to contribute to the measured low temperature rubber loss maximum.

The difference in loss relaxation area between the amorphous C2C3 n~bers (including Vistalon 5600) and that of the amorphous phases of the semi-crystalline C2C3 rubbers is striking, it might be (one of) the reason(s) for the differences in toughening efficiency between these rubbers. Futhermore, the amorphous phase loss areas of the in-situ

i00

Figure 4.3 Loss modulus of PP filled with 15%wt. of talc as a function of the reciprocal absolute temperature

7O

65

60

L'q E Z

55

r-- 5 0 (_9.-- v

E 0

~ - 4 5

v

E I,b I/) 0 ..J

40

35

30

25

20 2.40

+ +/ t'

I

,+,+

il + +

I I + +

I t

4- i+,++ + + /

+

/ f

f /

/ /

/

I I

2.83 3.26

+

f

+

~---I- - ' -+ - "+ - - + ' - - + --'4

,.I I I I

3 . 6 9 4 . 1 1 4 . 5 4 4 . 9 7 5 . 4 0

1 0 0 0 / T , K ,, - 1

i01

Figure 4.4 Results of Figure 4 .3 after subtraction of the background losses

4 5

4 0

3 5

0,1 E Z 3O

(9 ~ 25 r"

(b 0 " - -

. . . _ , , , . . . . .

0 " E (/) u) 0

_.l

15

10

5

0

/ + /

+

~ I -I-

+ ,~ i.% / , ' / , ' , \

I / I 2~ .... I " I I

2.40 2.83 3 .26 3 .69 4.11 4 .54 4.97 5 .40

1 0 0 0 I T , K ( - 1 )

0 0 0 I v

corr.

los

s mo

dulus

(G

"), N

.m2

0 PO

L0

Lo-

",4

4~

O

0 ~

I "'

I

'

,-,

~-I-..,

..~.i_

-.

) ~•

, --

~~+

~+

~+

~

'--.

,~+~

+~

+._~

+_.j:

l:

- §

_......

4...--

-"+ "~

+~

(Milli

ons)

O7

0 O7

0 ~

0 C~

I '

I '

I '

I '

I '

I '

C)~

c=

o'm

b~

0"3 ~-

~a

,..i.

X

~'~-

~o I _~.

~r

103

polymerised systems fit fairly well with the data of the mechanically blended semi-crystalline rubbers.

Table 4.1 Amorphous rubber fractions versus loss relaxation areas (DMA analysis of PP/C2C3 rubber blends)

r

I amorphous rubber fraction

i Nordel 1500 0.085 0.128 o.~vo

,, ,

I Nordel 3391 0.088 0.132

...... 0.Iv6 , ,

sys. A 0. 019 sys. B 0. 022 sys. C 0. 154 sys. D 0. 083

_ ~ - J L If ,

Ill Ill I

loss relaxation area Nm-2. K-I

,

2.2 E3 3.4 E3 4.5 E3

f

2.3 E3 3.9 E3 4.8 E3

0.7 E3 0.2 E3 4.0 E3 1.9 E3

amorphous rubber fraction

,,

Vistalon 404 0.I0 0.15 0.20

Vistalon 4608 0.10 0.15 0.20

..... ,, ,,,,

Vistalon 5600 0.096 0.144 0.192

, ,, , ,, ,

I I r ll I , ,I'

loss I relaxation area Nm-2. K-1

4.1 E3 5.5 E3 6.4 E3

3.6 E3 5.3 E3 6.7 E3

,,

4.0 E3 5.0 E3 7.3 E3

a. Nordel 1500 C2C3 rubber with 15 %wt crystallinity i.e. 10 %wt. Nordel 1500 in PP results in an amorphous rubber fraction of 0.085.

b. Nordel 3391 C2C3 rubber with 12 %wt. crystallinity. c. Vistalon 5600 C2C3 rubber with 4 %wt. crystallinity. d. Vistalon 404/4608 crystallinity < 0.5 %wt. e. Systems A, B, C and D are experimental, in-situ polymerised

TPPs, the amorphous rubber content of these rubber phases were determined by NMR on the extracts of the 'hot xylene solubles' (HOXS) procedure; the reported loss area is the difference between TPP loss area and the HOXS residu loss area.

8

104

Figure 4.6 Loss relaxation area versus amorphous rubber content

+ sere~- A amorph, o /n - -s i tu c. C 2 C C 2 C 3 s y s t e m s

I

O4 I E Z cO I..U X

r k,_ cO

c 0

X

o~

0 __1

7

6

5

4

3

2

0 0.00

A

/ o t I , I . i . I , . I , . .

0.04 0.08 O. 12 O. 16 0.20

Amorphous rubber f ract ion

105

4.1.3 Tu-value determination of aqed. riuid pU foams bv DMA Insulated pipes used for district-heating Systems consist of a steel inner pipe, a rigid polyurethane (PU) foam insulation layer and a PE pipe on the outer side. This foam should be capable to withstand prolonged contact with the hot steel surface of the inner pipe, retaining thereby its original high thermal insulation value.

Rigid PU foams are the reaction product of an excess of an isocyanate, a polyol and some water resulting in a polymer network with a high crosslink density. The excess of isocyanate used is indicated by the MDI index i.e an index of 135 means an excess of isocyanate of 35 %. The water added also reacts with the isocyanate under formation of C02 which acts as an 'internal' blowing agent. An 'external' blowing agent, evaporating by the reaction heat, is usually added also to obtain the desired foam density.

New PU foam formulations are anaerobically aged to test the heat resistance of the foams. The ageing process is monitored by measuring foam properties like weight, density, strength and thermal conductivity as a function of ageing time/ temperature. The Tg-value of such a foam was thought to be also an interesting property to monitor the ageing process in view of its close relation with the chemical structure of the foam.

Straightforward DSC Tg-value determinations on such foam samples failed, which was ascribed to too low sample weights. Cyrogenic milling of foam samples, followed by pressing the powder obtained into sample pills increased the sample weights from about 3 mg. to about 15 mg. The sensitivity of the DSC still proved insufficient to produce reliable Tg-values of these highly crosslinked systems.

Subsequently, a foam strip of 8 x 12 x 2 ~ was carefully clamped in a Polymer Laboratories DMA system and measured using a frequency of 10 Hz. and a heating rate of 1"C/minute. Figure 4.7 shows the result of such a measurement. The decrease of the Youngs modulus at temperatures > 120oC, accompanied by clear maxima of the loss modulus and the tan 8 indicate that the glass-rubber transition of such a foam is easily measured by DMA. The loss modulus (E") maximum temperature was chosen to indicate the Tg-value of these rigid foam systems.

The Tg-value development during anaerobic ageing of six differently formulated foam systems was then measured. Formulation 2 (MDI index 115) and formulation 1 (MDI index 135) show the effect of a MDI index decrease from the standard value of 135 to 115. These two systems also containted some glycerol next to the standard polyol. Formulation 3 is equal to formulation 1 but contains no glycerol. The formulations 3 to 6 (MDI index 135) were added to show the effect of a partial replacement of isocyanate by respectively 10% (4), 15% (5) and 20% (6) of a liquid epoxy resin.

L o g E " Figure 4 . 7 L o g E " r Results of the DMA analysis on the rigid PU foam sample (Pa )

i i i in II I I I IN I nn _ I I _ II ? .4

8 . 8

8 . 2

5 . 8

Tan d e l l ; a - E ' / E "

I I I I � 9 1 4 9 I I l ~ � 9

X �9 . 4 . _ ~ X 4-

T

�9 §

�9 X 4-

�9 X §

.#.~+#" �9 x +

~ X , ~ �9

m �9 m ~

i ~

� 9 s � 9

__..,,�9149 - l . _ . _ l l a l l a o a o o � 9 1 4 9 I l l l m l d l o l N l m q m ' -

x X

§ X

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§ .

4-

+ i §

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8

5 .9

5 . 0

5.?

5 .8

IZIMTFI

IB Hz STRRZH -,x4 L degC,,'ml n -LOGk- 3.?? SZNGL.s CFICr 2x t 2 x ? . l h m

CLFII4~ H/C F ' Z I . ~ t B a 2 1 4

B Y H I ~ ON O0O511O

I-= 0

107

The results of DMA Tg-value determinations on these foam samples during ageing at 160oC, are listed in Table 4.2. These results show three different stages in the ageing process of these foams: - first, the Tg-value increases i.e. the fresh, non-aged foam

is not completely cured and the crosslinking reaction(s) proceed due to the relative high ageing temperature,

- then, the cure processes are completed and the Tg-value remains constant for a certain period of time and, finally,

- the thermal decomposition process starts and results in a decrease of the Tg-value as a function of the ageing time.

The three Tg-value/ageing time curves in Figure 4.8 show these three stages.

The Tg-value of these foams with a more or less standard formulation i.e. the systems I, 2 and 3 is about 160~ after preparation. The ageing time at 160~ after which these six foam systems (during the third stage) again reach this Tg- value is used as criterium to compare the (long term) thermal stability of these foams. These times are plotted as a function of maximum Tg-value of the foams in Figure 4.8 (insert). The partial replacement of MDI by epoxy resin clearly increases the maximum Tg-value of the foams and thus their long term thermal stability. A MDI index increase from 115 (formulation 2) to 135 (formulation I) results in a Tg(max.) increase of about 20oC and this also improves the long term thermal stability.

Additional Tg-value determinations on foams aged at different temperatures proved subsequently, to supply excellent data for foam life-time predictions.

Table 4.2 Results of DMA Tg-value determinations on aged PU foams

form. number

1

2

3

4

5

6

Tg before

. .

1 6 1

160 ,, ,

160 r

180

1 8 3 . . . .

1 9 0 ......... ~ i,, i i

[ ~ p~ )i I

189 . . . . . . . . . . .

176 ..

183

193 . . . . . . . .

197

199

Tc

1{ 0"C "C

195 ,

175

186

196

198

202

T,g, 3 m,3nth 1,50"C 0C

190

158 , .... ,

184

195

198 ,,, ,,,

201

~79 th 6 TllC I

. .

184

142 , , , , ,

155 , , , i LI , , _ t

197 i i

203 . . . . . . .

Tg, 9 month 160oc oc

I III

141

184

> (.Q

(1)

(.Q

,-1.

3 (1)

m

,,,,_1

,,,

�9

(b

3 O ::3

::3"

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4~

03

�9

IX)

,.,,.,

Ik

03

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I I

Tg-v

alue

(E

"

k.)

(.,0

rj'l

(.71

I '

I

max

.), d

eg.

C

~ ~

--I.

~ ~

IX3

+ o o

+ I>

O

'~

Age

ing

time

at 1

60~

mon

th

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t4

~ U~

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r

ID

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.

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- .

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=,

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=,,,,

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- t~

r t~

ID.

:3

~3

3 0 0 o

109

4.2 Mechanical measurements at ultrasonic frequencies

4.2.1 Introduction The term ultrasonics is used to describe mechanical waves propagated in gases, liquids and solids at frequencies above the upper limit for the human ear i.e. above 16,000 Hz. The characteristics of these waves are related to the mechanical properties of the medium through which they pass. Ultrasonics can be used, therefore, to investigate these properties. From a practical point of view one can consider two separate cases, one where the wavelength is much greater than and the other where the wavelength is much shorter than the sample thickness. The first case is usually called acoustic elastometry (frequencies ranging from 20 kHz. to 200 kHz.) and the second is called ultrasonic elastometry (typical frequencies 0.5 MHz. to 5 MHz.).

When ultrasonic waves pass through a medium, the particles of that medium start to vibrate at ultra-audible frequencies. Part of the energy of the vibrating particle is transmitted to neighbouring particles. Because subsequent particles start their vibration one after the other with a slight time delay, the vibrational motion travels with a finite velocity c known as the wave velocity. The phenomenon is described as wave motion. Two kind of waves can be propagated in infinitely solid media- longitudinal waves and transversal waves. Longitudinal waves can be propagated in all types of media, the particles of the medium vibrate then in the direction of the propagation. Transverse or shear waves need shear elasticity and can therefore only be propagated in solids while the particles of the medium vibrate in a direction perpendicular to that of the propagation.

The ultrasonic velocity and absorption (measured to examine the elastic properties of a material) can be determined using different experimental methods. A describtion of these methods is given by Philipczynski et al. [8]. The immersion technique in combination with the pulse propagation technique is commonly used to investigate polymeric systems [9, i0]. In the immersion technique, the sample, transmitter and receiver are all immersed in a liquid. Ultrasonic pulses are sent from the transmitter to the receiver both with and without the sample in the path of the sound beam. Longitudinal waves are developed in the sample when the sample is held perpendicular to the path of the sound beam. If the sample is held at an certain angle to the sound beam both !onuitudional and shear

. . . . . . . . . . .

waves are generated in the sample. The longitudinal waves are totally reflected and only the shear waves are propagated if the angle at which the sample is held is greater then the so- called critical angle.

The immersion of the system in a proper liquid ensures a good coupling between transmitter~sample~receiver and promotes" the temperature control of the system. When a wave is incident normally to the boundary between two media both transmission and reflection occur. ~"ne reflection coefficient is high for

ii0

waves passing from a solid or a liquid to a gas i.e. ultrasonic waves propagated in solids/liquids do not penetrate in the surrounding air.

With the sample perpendicular to the ultrasonic beam and the basic values listed below known or measured: L = the sample thickness, m 1 - the transducer distance, m tl - the pulse propagation time without the sample, s A1 - the pulse amplitude in the liquid, V t2 - the pulse propagation time in the liquid/sample/liquid, s A2 = the pulse amplitude in the liquid/sample/liquid, V pl = the density of the liquid, kg/m3 p2 = the sample density, kg/m3 the ultrasonic properties can be calculated [8, I0].

The propagation speed of the waves in the liquid v(liq) is:

v(liq) = i/tl, m/s 4.13

tl- I/v(liq) and t2 = (i-L)/v(liq) + L/v(1) hence,

tl - t2 - L/v(liq) - L/v(1), the longitudinal propagation

speed of the waves in the sample v(1) is:

v(1) = [v(liq).L]/[L - v(liq). (tl - t2)] 4.14

The propagation speed of the shear waves in the sample v(s) is-

v(s) - v(liq). [(cosu - v(liq) . (t2'-tl)/L)' + sin'a]*" 4.15

where u is the angle of incidence between sample and sound beam and t2' is the pulse propagation time through the liquid/ sample/liquid in the shear mode. For the waves travelling from the liquid into the sample holds:

sinu/v(liq) - sin~/v(1)

The angle u is critical if the sin~(1) = 1, hence

sinu(c) = v(liq)/v(1)

4.16

v(1) and v(s) can be used to calculate the elastic constants G, B and E in N/m2 and the Poisson ratio ~-

G = [v(s)]'.p2

B = [v(1)]'.p2 - 4G/3

E = 3G/(I + G/3B)

= 0.5 - E/6B

4.18

4.19

4.20

4.21

4.17

iii

Amplitude A2 is smaller than amplitude A1 due to internal friction effects in the sample i.e. absorption takes place (the liquid absorption << rubber absorption is neglected). There is also signal loss due to reflection losses. The reflection losses of the two rubber surfaces can be calculated according to [i0] :

Lr - Ln[(Zl + Z2)~/(4.ZI.Z2)] 4.22

where Z is the acoustic impedance i.e. Zl = pl.v(liq) and Z2 - p2.v(1).

The total signal loss is expressed as-

Li = Ln AI/A2 4.23

The attenuation coefficient of the sample for longitudinal waves u(1) is then calculated in Nepers/m by-

G(1) - (Li - Lr)/L 4.24

The absorption losses can also be expressed in an ultra-sonic loss factor according to [9] �9

Loss factor = [u(1).v(1)]/~.f 4.25

4.2.2 The ultrasonic measuring ..... ~quipment An ultrasonic measuring system was developed to measure the attenuation coefficient/loss factor of car-tyre rubber samples as a function of the temperature between about -40"C and 150"C. These measurements were performed on sample disks of vulcanised rubber having a diameter of 60 mm. and a thickness of 13 mm.

The sample holder system used contains six sample apertures. Five samples are maximally placed at the same time in this holder to keep one aperture free for the reference measurement, see Figure 4.9. This whole sample holder system is lifted into a special thermostat bath provided with a liquid nitrogen cooling coil. This cooling possibility extends the lower temperature limit of these measurements from 20~ to about -50"C. The bath is filled with a mixture of water/ethylene glycol (1/1) for measurements between -50oC and 80~ Silicone oil (100 cS.) is used as medium for measurements between 0oC and 200"C. The sample temperature is measured by a platinum resistance thermometer, placed as close as possible to the sample in the ultrasonic beam.

The sample holder can turned around in the YZ-plane (the transducers axis being the X-axis) to bring sample after sample into the ultra-sonic beam. The whole sample holder system can be rotated in the XY-plane to change the angle of incidence between 90" for longitudinal wave measurements, to 0 ~ to perform shear wave measurements. Both transducer arms are moveable in the X-direction to change the

112

Figure 4.9: US total immersion sample holder system

113

transmitter/receiver distance. Besides, a micrometer controlled movement of both transducers in the Y-direction is possible. This offers the possibility to correct for deviation of the ultrasonic beam during shear wave measurements.

One inch immersion transducers (Panametrics) are used with resonance frequencies of respectively 0.48 MHz., 0.70 MHz., 0.90 MHz. and 4.0 MHz. These transducers can be used for measurements between -50"C and 80~ The two transducers shown in Figure 4.9 are special high temperature transducers (resonance frequency 1.0 MHz.) which can be used up to 200~

Figure 4.10 shows the block diagram of the electronic setup. A pulse generator produces a 200 Hz. rectangular voltage which negative going side triggers trigger unit one (T1). The five #s. pulse produced by T1 is used to trigger the timebase of the oscilloscope and to start trigger unit two (T2). The timebase of the oscilloscope thus starts 5 #s. before the ultra sonic pulse begins, to display both the transmitting and the receiving pulses completely on the scope. The T2 pulse is adjustable between 1 and 80 ~s. A pulse length of 25 ~s. is commonly used. The T2 pulse triggers the HP3312A function generator which produces a sine wave 'filled' pulse of 25 #s. This pulsed signal is amplified to a signal of about 200 Vpp. and send to the transmitting transducer. T2 also starts the HP5327B counter.

The receiving transmitter produces a signal of about 28 Vpp under standard conditions (temperature 20~ frequency 0.90 MHz., transducer distance 0.10 m. and a water/ethylene glycol medium). The amplitude of the signal received is read-out via scope channel Y1. The adjustable level detector permits a manual choice of the trigger moment on the starting-flank of the receiver pulse, even if there is some signal distorsion due to high absorption levels of rubber samples going through their glass-rubber transition regions. The detector triggers a one-shot circuit which produces the counter's stop pulse. The pulse travelling time (counter read-out) and the receiver transducer output signal (oscilloscope read-out) are measured with and without a sample in the ultrasonic beam and used subsequently to calculate the ultrasonic properties.

4.2.3 Results of ultrasonic measurements on car-tvre ~lhbers. The physical properties of a car-tyre rubber compound are a compromise between a number of often opposite demands, see 1.2.1. Two of such demands are the holding behaviour on wet roads, the wet grip (WG-value) and the rolling resistance (RR- value). The contribution of the rubber phase to these two properties seems to be clear: increasing dynamic mechanical losses are in general causing an increase of both properties. This is good for the WG-value which has to be high for safety reasons. The RR-vaiue, however, has to be as low as possible for economical reasons.

200 Hz. -| -]

I OSC.

114

T E K T R O N I X 4 6 8

I _ T1 J L I

_ i

, ,,, ,= ,,, , ,,

scope trigger digital storage

__ _ scope

I Y2 Yl I "

T 2 . . . . samp 1 " '

I ii_. �9 ~ , - .

f u n c t ~ on RF amp I ~ f ~ er g e n e r a t o r

,, ,, , , -

. . . . . . . . . . HP 5327B timer/counter

start stop

200 Hz. ,

T1

T2

!

5

, , H . p s .

I I I

" ~ 2 5 ~s.

_ 1 - I ! ! !

I I I

F [ _

Figure 4.10 US electronic diagram

115

However, there seem to be possibilities for a compromise. The RR-value is related mainly to the dynamic mechanical losses at frequencies of about 10 Hz. in the temperature region 30~ - 60~ The WG-value, on the other hand, is the sum of a macro hysteresis and a micro hysteresis effect. The dynamic losses in the glass-rubber transition region of the rubber at frequencies of about i0 Hz. seem to correspond with the extent of the macro hysteresis contribution, whereas the dynamic losses in the high frequency region (40 kHz. up to 10 MHz.) seem to correspond with the extent of the micro hysteresis effect. This might offer possibilities to increase the micro hysteresis i.e. the WG-value by influencing the rubber structure at an equal or only slightly increased RR level.

The dynamic mechanical losses of a series of six (emulsion) styrene butadiene rubber (SBR), (solution) styrene butadiene rubber (SSBR) and butadiene rubber(BR) samples with a known (laboratory) WG-rating were measured in their glass-rubber transition region. This, to investigate if a high frequency loss/temperature region, especially related to the WG-rating of these samples, could be specified. Figure 4.11 shows four of the experimental curves, the loss factor values being calculated according to equation 4.25. Subsequently, the correlation coefficients between the ultrasonic losses of these samples at a certain temperature and their WG-ratings were calculated with a temperature interval of five degrees. The results of these calculations are plotted in Figure 4.12. The curves in this figure show that the ultrasonic losses measured at 5~ MHz and at 10~ offer the highest correlation coefficient values. The WG-rating/ultrasonic loss (10~ relation is shown in Figure 4.13. The commonly used DSC Tg(onset)/WG-rating relation is also shown in the inserted figure. These data show that the correlation WG- rating/rubber property improves considerably if high frequency loss data are used instead of the commonly used (low frequency) DSC Tg-value data.

The mechanical properties of an experimental SSBR sample were measured between -30oC and 5oC, see Table 4.3. A data set used to calculate these properties according to the equations given in 4.2.1 is given below as an example:

Sample : black vulcanised SSBR (50 phr carbon black), Sample thickness: 0.0134 m., Frequency : 0.48 MHz., Temperature : - i0 o C, Medium : silicone oil,

v(liq) : 1189.0 m/s, v(1) : 2404.5 m/s,

The critical angle necessary to measure the shear velocity is then according to equation 4.17- 29.6 ~

0

I 0

0 I ID

0

--I

0

o

0 0 | 0

0 0

0 0

0 0

0 b

b b

b '~

"~

~

x;'"

,~'"

,"

~,

'

' "~

'~

''

' "

"X~

~Los

s fa

ctor

(0.9

4 MHz

) "1

>,,~

- +

+

o

- /

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~~

J"

/ ~

~§

\

- .

~ ~

__J-

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.

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.

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z

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/ 0

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~+

o -

-~0

~ t~

o

�9

~3

0 0

r-

"TI

o~

"

+

~~

~

N

3 m

c

0 _~

.

0 0

o~

:3

.-h

W

0 0 13

"

§ =

o ~

0

(I)

(1) D) 3 ,I

o~

E)

W

0 0 0

0 L0

0

0 0

0 0

......

, ~

........

. ~ ~

~,

.

..

..

,

Corr

elat

ion

coef

ficie

nt

~ !>

~

~ +'

~x

~!>

~

j+

_ ~ _

1_~

EI"~

+ t>

0 /

0 / 0 / 0

b 0

C

C

I> E

= N

f~

0 iI

0 m

0

b N

Q < C

118

Figure 4.13 Correlation between ultrasonic loss at 10~ and 0.9 MHz and the W.G. rating

0 . 1 6

0 .14

0 .12

0 . 1 0

-31o

0 . 0 6

0 .04

0 .02

0 . 0 0 G

7 0

N 1-

0 .08

0 0

0 m

o c 0

D

+

+

r , ,2 - 0 . 9 9 5

Correlation between the DSC onset Tg-value and the W..Q. rating

r ,2 = 0.924 /

. / +

- ' 4 O

0 o

6 --~o m

i

IU

- 7 0 r.. O

I--- .-eo O 03 -9o Q

t- - 1 1 0 �9 , �9 , , , , , �9

7 0 8 0 9 0 1 0 0 1 1 0 t20

Locked wheel Wet Grip rating

8 0 9 0 100 110 120 130 140 15r

Locked wheel Wet Grip rating

119

This experimental situation is drawn in Figure 4.14. The deviation of the ultra-sonic beam from the 'main' X-axis, due to the shear waves is clearly visible.

The micrometer controlled movement option of the receiver in the Y-direction is used to move the receiver in such a way that the receiver signal again reaches its maximum value. V(s) is calculated, subsequently, according to equation 4.15.

v(s) �9 948.5 m/s,

Using v(1), v(s) and the rubber specific weight at -10~ (945.3 kg/m3) the elastic properties can be calculated using the equations 4.18 - 4.21-

G modulus : 8.5E8 N/m2, B modulus �9 4.3E9 N/m2, E modulus : 2.4E9 N/m2, Poisson ratio : 0.41

The results are plotted as a function of the temperature in Figure 4.15. Measurements like this are time consuming mainly due to the rather long times necessary to reach constant temperature conditions. These results illustrate however the advantage of this technique, measurement of all three moduli and the Poisson ratio during a single experiment.

Table 4.3 Results of moduli measurements at 0.48 MHz. on SSBR rubber

, , , , ,,

temp. , "C

5

0 .,

-5 ,

-10

-20

-30

, . ,,, . _ , , ,

B m o d u l u s xE9 N/m2

3.1 , ,,

3.7

E modulus xE9 N/m2

' , " I - I

G modulus xE8

3 . 7 t , , ,

5 . 2 . . . .

6 . 8

i.I

1.5 , ,, ,,

4.1 1.9 +

L . . . . . . . ,,

4.3 2.4 -- , i ,

4.8 2.9 ......

5.2 3.2 _ +

,,,, .,

Poisson ratio

0.44 , , ,,

0.43 ,,

0.42 L �9 ,, ,

8.5 0.41 �9 , , , t ,

I0.4 0.40 _ ,, , ,-,,

11.5 0.40 " ,,,, _ + I,~,,

* medium: silicone oil

/ Rubber samp l e

R e c e i v e r

Figure

Shear

v(s)

4.14

velocity

I I 4III |

/ vci)

measurement

II(i) - 29.6 I i(o. I) = 90 e

i(o.s) " 23.2 *

I

~ t - I Transmitter I

S a m p l e : s o l u t i o n SIBR T e m p e r a t u r e : - 1 0 C M e d i u m �9 s i l i c o n e oi l T r a n s d u c e r s : P a n a r n e t r i c s 0 . 5 k44z.

0

121

Figure 4.15 Elastic moduli of solution SBR (measuring frequency 0.48 MHz)

+ K A E 0 G + ~u m o d m o d rood ra rio

l e + 10 . . . . . 0 . 5 1

Cq E Z

. . . . . .

0

l e + 0 9

l e + 0 8

- 3 5

O ~ ~ 0

j ' - o 0 § \ / / X

4- 0

+

J + . - -

i I , , I j I J _ I

- 2 5 - 1 5 - 5 5

- 0 . 4 9

i

0 . 4 7

- O . 4 5

- 0 . 4 3

- Q . 4 1

0 . 3 9

15

T e m p e r a t u r e , deg. C

122

References

1. B. Wunderlich- Thermal Analysis, Academic Press Inc., New York, 1990.

2. K. Schmieder and K. Wolf, Kolloid Z., 134, (1953), p. 149. 3. L.E. Nielsen and R. Buckdahl, J. Appl. Phys., 21, (1950),

p. 607. 4. J. Heijboer, Plastica, 19, (1966), p. 11. 5. J.M. Charlesworth, J. of Mat. Sc., 2~, (1993), p. 399-

404. 6. J. Heijboer, Mechanical Properties Of Glassy Polymers

Containing Saturated Rings, Thesis TU Delft, (1972) . 7. J.J. Fay, D.A. Thomas and L.H. Sperling, J. Appl. Pol.

Sc., ~/, (1991), p. 1617. 8. L. Filipczynski et.al.: Ultrasonic Methods of Testing

Materials, Butterworths, London, 1966. 9. H.A. Waterman, Kol. Zeitschrift & Zeitschrift fur Pol.,

192, 0kt. 1963, p. I. 10. B. Hartmann and J. Jarzynski, J. Acoust. Soc., ~_~, 5,

(1974), p. 1469.

THERMO-ELECTROMETRY

CHAPTER 5

123

CHAPTER 5" THERMO-ELECTROMETRY

5.1 The DC and AC properties of polymers

5.1.1 Introduction Thermo-electrometry is a group of thermo-analytical techniques in which an electrical property of a sample is monitored as a function of the temperature or time. An electrical property is seen as the response of a polymer when an electric field is applied to it. In contrast to metals, where electronic conduction is the only response to an electrical field, polymers may respond in different ways. A review of the different possibilities is given recently by C.C. Ku and R. Liepins in their "Electrical properties of polymers; chemical principles [I]. Ku and Liepins separate the response of polymers to an electric field into two main parts-

i. The dielectric properties and 2. the bulk conductive properties.

These two parts are subsequently split up into four fundamental parameters, characterising the dielectric properties of polymers-

IA. The dielectric constant, representing polarisation and

lB. the dielectric loss angle, representing relaxation phenomena.

2A. The conductivity, representing the electrical conduction and

2B. the dielectric strength, representing breakdown phenomena.

In addition to these four fundamental parameters, special electrical properties are recognised like: piezo-, pyro-, ferro- and tribo-electricity and photo voltaic/conducting properties. The contribution in this chapter will be limited to three of the four fundamental parameters: AC measurements (IA/IB) and DC measurements (2A). Besides, attention will be given to a kind of combination of AC and DC measurements- the thermally stimulated discharge (TSD) analysis technique. An analysis technique used to detect relaxation phenomena in organic and anorganic materials.

124

5.1-2 DC properties of polymers Polymers, although often considered and used as electrical insulators, are never ideal insulating materials. There are always some charges present which are free to move in an electrical field (E). For polymeric systems these charges are mainly ionic species. The current density due to the presence of an electrical field may be described in general by-

j (i) - n(i).z(i).v(i) 5.1

where: n(i) = the number of charge carriers of type i per unit of volume,

z(i) = the charge on each carrier and, v(i) = the mean velocity of this type of carrier.

N(i) and v(i), and - if different ions are present - z(i) can be influenced by common variables such as temperature, pressure and electrical field strength. In contrast to good conductors like metals, the specific volume resitivity [E/j(i)] of polymers is-

- field-strength dependent, - decreasing with increasing temperature and, - increasing with increasing pressure.

A rectangular sheet of polymer between on both sides a conducting plate forms a capacitor. Such a capacitor can be used to measure the polymer's volume resistivity. Application of a voltage (U) over both conducting plates results, after a certain period of time, in a constant current flowing through the polymer sheet I(dc). Then holds:

- (Rx.A)/d and Rx - U/I (dc) 5.2

where" G = the specific volume resistivity, Ohm.m A = the conducting plate area, m2 d = distance between the conducting plates, m

Such an experiment shows, however, that it takes time to reach this constant I(dc) value. This is caused by a polymer effect, the polarisability.

Application of a voltage (U) over both conducting plates without the polymer sheet inserted, results immediately in charges +Q and -Q on the plates. Charge and voltage are linearly related according to.

Q = Co.U = ~o. (A/d) .U 5.3

where: Q = the charge in Coulombs and, co = the dielectric constant in vacuum,

(8.85E-12 Farad/m), Co = thr capacitance of the parallel plate system.

An increase of Q to Q' at the same applied voltage is measured some time after the polymer sheet is inserted between the charged conducting plates. For Q' holds-

125

Q' = Cx.U = ~o.cr. (A/d).U = Er. Co.U 5.4

where. ~r = the (static) relative dielectric constant of the polymer

This increase in charge is, for non-polar materials, due to two effects [i]-

I. Electronic polarisation, due to the displacement of the electrons relative to the atomic nuclei. The characteristic time for this contribution, de, is below E-16 s.

2. Atomic _ Dolarisation, associated with the displacement of the atomic nuclei relative to another, da, with characteristic times of the order of E-13 s.

With polar materials like many polymers, a third contribution arises from the orientation of uncompensated dipoles in the electric field. These dipoles can be permanent dipoles which exist in the absence of an applied electrical field, and are due to differences in electronegativity of the bonded atoms (for example the carbonyl bond C=O). They can also be induced dipoles, created by the applied electrical field which causes redistribution of electrons shared between bonded atoms with similar electronegativity. This dipole orientation is counteracted by the brownian motion of the groups carring the uncompensated dipoles. Consequently, the characteristic time for these effects is strongly dependent on the temperature. Its lower limit is of the order of E-II s. for small molecules and, often, of the order of E-2 to E+4 s. for polymers at room temperature. Since this Qrien~ation polarisability, do, involves physical movement of parts of the macromolecules it is not surprising that, apart from the temperature, other factors which determine the molecular mobility affect it. Thus, the total polarisability is given by-

dT = de + da + do

This polarisability now causes that the charge increase from Q to Q' needs a clearly measurable amount of time. Thus, a charging current (Ic) decreasing with time, superimposed on I(dc), flows if a voltage (U) is applied over the plate/ polymer/plate capacitor. The proper I (dc) value necessary to calculate the specific volume resistivity, can be measured only if Ic has become zero.

The specific volume resistivity/temperature or electrical field-strength relation is measured during what usually are called DC measurements. A specific method was developed to determine the volume resistivity of a polymer as a function of the temperature and/or the electrical field strength [2] without the need to wait until the Ic current component has become zero. This method is schematically given in Figure 5.1.

126

SAMPLE r-"R".-,

H I ~ ! L I I i C x I

U i . . . . . I Io] , , ,

DETERMINATION OF ];O SAMPLE ',-,;-~G.

_H, j ~ Z

TIME

is§ - j TIME CHARGING

SAMPLE i ~:::::::::~ i L

- - I I - - - I ~ ' Y

U

DISCHARGING

Current f lows during volume resistivity measurements

Figure 5.1

127

A metal plate/polymer/metal plate combination often causes an extra current contribution (Io) due to electro-chemical processes and/or contact potential differences. Therefore, the sample/ electrode system is, therefore, first short-circuited via the electro-meter. The value of this Io current (§ is monitored until it has become constant. Then the sample system is connected with the voltage supply and during a certain time a charging current Ic is measured which is the summation of the pure capacitive charging current, the dc current and the Io contribution. During the third step, the sample system is again short-circuited via the electrometer. Now, a discharge current I(d) is measured which is the summation of the pure capacitive discharge current and the Io contribution.

I (dc) can now be calculated because for equal charge/discharge times holds-

I' (dc) = I(c,t) + I(d,t) [I(d,t) has a negative sign] 5.5

This value has still to be corrected for the Io contribution i.e.-

I(dc) = I(c,t) + I(d,t) - 2.Io 5.6

I (dc) is inserted in equation 5.2 to calculate the specific volume resistivity at that temperature and electrical field strength. This measuring procedure is completely automated, see 5.1.4. The volume resistivity determination is described in the ASTM D257 (US), BS 202A (UK), DIN 53596/51953 (BRD) and in IS0 93 [3]. The time dependency of Ic is mentioned in all these methods. Besides, all methods indicate that I(dc) is determined usually after a standard charging time of 60 seconds. It will be clear that if one assumes I(dc) = Ic(60 seconds), this current nearly always contains a time dependent charging current contribution.

128

5.1.3 AC properties of polymers The voltage applied to the conducting plate/polymer/conducting plate system described in 5.1.2 i.e. Cx and Rx in a parallel circuit, is sinusoidal during AC measurements :

U = Uo.sinet 5.7

The result is a current through Cx given by:

I' = dQ'/dt - Cx.Uo.~.cos~t 5.8

and an additional current through Rx of:

I' ' = [Uo/Rx] .sinet 5.9

I(total) = I' + I'' = Uo.[Cx.~.coset + (I/Rx).sin~t| 5.10

According to 5.4: Cx - or.Co 5 .II

thus: I(total) = Uo.Co.e.[~r.coswt + (i/Rx.w.Co).sin~t] 5.12

Now, two frequency dependent, material properties can be defined as:

c'r = or, c''r = I/Rx.~.Co

c'r = the out-phase, capacitive component and is called relative dielectric constant;

~''r = the in-phase, resistive component and is called the relative dielectric loss factor.

~'r and ~''r are the components of a complex quantity called the relative permittivity: E*r.

The phase difference between I (total) and U is the loss angle described by the tangent ~ or dissipation factor-

tangent ~ - ~''r/c'r = 1/Rx.w.Cx 5.13

The polarisation effects mentioned in 5.1.2 are characterised by their time of response to an applied field or their "relaxation" time. In polymers in particular, there is a distribution of relaxation times rather than a single characteristic relaxation time, usually with a fairly high maximum. Height, location of the maximum and width of the distribution are often used as characteristics.

During AC measurements the frequency determines the dipole response: if the relaxation time T and the frequency ~ are such that

e.T >> 1

the dipoles are not capable of following the field and do not contribute to the polarisation and thus the dielectric constant. If, however

129

~ ) . T < < 1

they fully contribute. In both limits w ~ w and e ~ 0 no energy is spent due to the lack of either dipole movement or opposing forces. In between these frequency limits energy appears to be spent, in other words, there is a form of conductance present which by its nature is frequency- dependent. If

W.T ~ 1

there is finite movement and finite opposing force, resulting in a maximum in dissipated power. This corresponds to a maximum in the phase lag between field and polarisation.

The magnitude of e. T determines the response of a dielectric. To study this response either w or 7 may be varied. A wide variation of e means, experimentally, that different apparatuses have to be used. T, however, is easily and widely varied by varying the temperature. In polymers often more than one type of relaxation mechanism is encountered. Each has a specific average relaxation time, distribution of relaxation times and temperature susceptibility. The temperature susceptibility can often be described by an Arrhenius factor [i] so that:

T = ~o.exp(E*/RT) 5.14

where. To = the characteristic relaxation time, E* -- the energy of activation, R = the gas constant and T = the absolute temperature.

Measuring the dielectric loss maxima as a function of the temperature at a number of discrete frequencies provides the data for an Arrhenius plot i.e. in(~) versus i/T(max.). According to equation 5.14 and ~(max.). T = I, an experimental activation energy value can be calculated from the slope of this curve-

E*/R = -d [In (Q) ] /d [i/T] 5.15

where: ~ = 2.~.f and f = the measuring frequency, Hz.

If the dielectric loss factor versus T at one fixed frequency is available, E* can be calculated from the area under the loss factor versus I/T plot. Substitution of 5.14 in Debye's equation for a system with a single relaxation time i.e.

E''r : ~.T.(E'o - E'.)/(I + eat ') 5.16

where: E'o = ~'r for f ~ 0 c'. = ~'r for f ~ w

gives-

130

E* = (C'o - E'.) .0.5.~r.R. [~E' 'r.d(i/T)) ^(-I) 5.17

The actual dielectric losses measured are larger than the dipole losses due to the contribution of the conduction losses i.e.

c''r(total) - c''r(dipole loss) + c''r(conductance loss)

The c' 'r(conductance loss) contribution can be calculated according to 5.12 �9

' ' r (conductance loss) = 1/[Rx. e. Col 5.18

where Rx calculated according to equation 5.2, is the value measured during DC experiments on the same sample. Equation 5.18 shows that the conduction loss is inversely proportional to the measuring frequency.

The discharge current (Id) measured during the DC experiments described in 5.1.2, corresponds to the dipolar losses while the I (dc) corresponds to the conductance losses according to:

e''r(total) = I(c,tl)/[w.Co.U]

= I(d, tl)/[u.Co.U] + I(dc)/[~.Co.U] 5.19

where- u = 2.~.fl and fl, the corresponding frequency, is according to Hamon [4] :

fl - 0.1/tl 5.20

The equations 5.19 and 5.20 offer the possibility to combine the loss data obtained with AC measurements and those obtained with DC measurements extending the lowest measuring frequency from about 10 Hz. to about 10"4/10 s Hz.

More extended discussions about the dielectric phenomena can be found in several textbooks [5, 6, 7] .

131

5.1.4 The AC and DC measurinQ svstem The automated volume resistivity/dielectric properties measuring system is schematically drawn in Figure 5.2. The whole system is controlled by the HP9000/300 computer which runs either the specific volume resistivity measuring program or the dielectric properties measuring program.

The polymer sample, clamped between a guarded electrode system is placed in the thermostat bath. This bath can be heated electrically and cooled by liquid nitrogen providing a measuring region from -150~ up to 300~ The bath inside is purged with nitrogen (about 200 ml/minute) to prevent sample oxidation at high temperatures and moisture condensation at low temperatures. A Cr/Al thermocouple measures the sample temperature. A second thermocouple, measuring the bath side- wall temperature, is connected with the temperature controller of the thermostat bath.

Activation of the computer gives a menu with the possibility to choose the 'Volume resist' program or two 'LCR' dielectric programs. The 'Volume resist' program is controlling the Keithely 617 electrometer, the Keithely 237 High Voltage Source (0 - i000 V) and the high voltage switching unit. This system performs completely automated the DC measuring procedure drawn in Figure 5.1 and described in 5.1.2. This measuring procedure can be performed at twenty different temperature steps, at maximally three different voltage levels each. The current is read-out by the electrometer and stored every thirty seconds during the Ic measurement and the Id measurement (both maximally 40 minutes). The temperature of the thermostat bath is brought to the next chosen measuring temperature after performing a complete DC measuring cyclus.

The 'LCR single' program measures the dielectric properties of one sample cell as a function of the temperature and frequency. The HP4284A LCR meter measures at frequencies between 20 Hz. and I MHz. The 'LCR multiple' program controlls an additional scanner which makes measurements on maximally five samples during one measuring scan possible.

The sample temperature is then measured either with a Cr/AI thermocouple or with a Ptl00 resistance thermometer placed in the high potential electrode of the measuring cell. Both sensors are connected with the sample temperature measuring unit which is connected with the HP9000/300 computer, see Figure 5.2. During the dielectric measurements the temperature of the thermostat bath is continuously increased or decreased usually at a rate of I~ minute. This in contrast with the volume resistivity measurements where the temperature is step- wise in- or decreased. The balancing time of the HP4284A LCR meter [seconds] permits however, the use of a temperature ramp. The sample temperature follows the thermostat bath temperature with a temperature lag of about 25~ The interval time chosen determines the temperature difference between two subsequent measurements.

132

SCR~NER I I I I I i I I I

H P 4 2 8 4 R p r e c ! �9 t o n LCR m e t e P

____t__---- SRMPLE TEHPERR- TURE

, _ . . . .

. , , . .

BRTH

TENPERR- TURE CONTROL- LER

, ,

HP9153B DISKDRIVE

HPBBBB/38B COMPUTER

i" J PRINTER

i" HP COLORPRO PLOTTER

HIGH VOLTRGE SWITCHING UNIT

KEITHLEY 237 HV SOURCE MERS.UNIT

i

. . . .

KEITHLEY 6 t ? ELECTRO-- METER

Figure 5.2 Measuring system (schematically) for determination of the non-destructive electrical properties of polymers

133

Samples, often received at a thickness between 3 and 5 mm., are usually machined to a thickness of 1 mm. Subsequently they are provided with a (vacuum evaporated) silver low- and high- potential electrode system. Sample disks with two different diameters i.e. II0 mm. and 53 mm. are used, with low potential electrode diameters respectively of 80 mm. and 20 mm. The samples with their silver electrodes are clamped between a spring loaded, guarded electrode systems from the sample cells. These samples and sample cells are used both for the volume resistivity measurements and for the dielectric measurements. The dielectric measurements are performed according to the instructions given by IEC 250 [8]. Sample cells for disk-shaped rigid samples, for high viscous/zltbbery and for liquid samples are also constructed according to the directions given in IEC 250.

5.1.5 AC and DC properties of a cured resin system _

Figure 5.3 shows an example of a dielectric measurement. The dielectric constant and the tangens delta of a liquid epoxy/ polyamide resin system measured between -50~ and 150"C at ten different frequencies between 60 Hz. and 1 MHz. during one experiment. The results of five selected frequencies are plotted in Figure 5.3. The resin system shows a broad glass- rubber transition region between about 30~ and 100~ indicated by the step-wise increase of the dielectric constant and the maximum in tangens delta. This sample was cured during 12 hours at room temperature followed by 6 hours at 100~ Heating up to 150~ during the dielectric measurements might have postcured the sample. The volume resistivity measurement was performed, therefore, on a second 'fresh' sample disk.

The volume resistivity was measured at 14 discrete temperatures ranging from -36oc up to I06~ Figure 5.4 shows the measured charge and discharge currents and I(dc,t) current calculated according to equation 5.6 as a function of time at a temperature of -4.7oc. The sample is then in its glassy state and Ic decreases as a function of time according to-

Ic m t.exp(-n) 5.21

Such a behaviour is also measured for PVC, PE and many other glassy and semi-crystalline polymers [9, I0, II]. The I(dc,t) values calculated make the determination of an average I(dc) value possible although the I (dc, t) values clearly scatter. This sample is just in its rubbery state at 19.6oc, under these measuring conditions. The scatter in the calculated I(dc,t) values has been practically vanished and the system approaches the ideal Ic/Id/I (dc) system described in 5.1.2, see Figure 5.5.

134

Liquid epoxyipolyamide resin

s.s~

5 . 0 o

11.S-

8 . 0 -

7 . 5 -

7. t l .

I ; . S .

S . 0 -

5 . 5 -

5 . e -

4 . $ -

4.8-

3,$-

- 50

o...._ o - / \ 01 O

0 k i l t . / | .g kH:. ~ o r | . l

~= / 7 ,,.o =_~=_.o " x ---, u ~ r V | / / / ~ = x t ' ~ 0 / ~,a IQID. e hH=. X _ ~

/ 7 / . . ~ o ' - - o ~ , , = ~ ' / / = " / " " ~ o ,+

/ / / +i+ 1.: P.,.,,., < / / / / - . - + - + I = o / , /

T ~ m r t a r R t l i m o C

' w ' ~ " ~ - " ' I I +

-30 - 1 0 1'0 :30 50 ;PO 90 110 130 1+0

Uquid epoxy/polyamide resin

8. I 4 - ~ / / 18.8 kH=. o - / / "10 1.= kH=. / / f z ~ \

0. tZ- C: / v =" t ~ / o At.0 ~ : .

o g.

J / i / , - ' -

O. 01; -

0.04 -

,o,_. o-~-:.~

a . ~ - " , ; -' ; 0 ' 5'0' - ' ' ' - 50 - 3 0 - 0 30 70 90

< /

T e m p e r a t u r e , ~ .... t I 110 130

Figure 5.3 Results of automated dielectric constant and Tan delta measurements

150

135

1 0 -10

1 0 -11

10 -12

Current, A Figure 5.4 Results of charge/discharge current measurements on a liquid epoxy/polyamide resin sample in its glassy state

T e m p e r a t , u r e = - 4 . 7 ~

+

. I | | |

N

IN)

- I ( d t s c h a r g e )

+

X ~ + ,

x\ \ \

I ( d c . )

l(charge) +

• \§ X + ,

\ + • \ \ + \

X +

X \ X

\ X

+

\ X

<3 <3

<~ <3 <3

Time, s

136

10 -9

1 0 -10

1 0 -11

G~

Current, A

Figure 5.5 Results of charge/discharge current measurements on a liquid epoxy/polyamide resin sample in its rubbery state

l ( d c . )

x

T e m p e r a t u r e = lg.6eC

I ( c h a r g e )

,+

<3 <3

x \

+~

" + ~ + . +

+ -I- + + . . + .~. +

<3 _...<=--<=-.~-<3 <3 <3<=-<~ --~

X - I ( d t s c h a r g e )

\ x

\ X \

x s x k X

X

X \

x

. i J i

N

I |

Time, s 9.--4

137

If the Ic-value after 60 seconds is used as 'I(dc)'-value, for both experiments shown in the Figures 5.4 and 5.5, this 'I (dc)' value contains a time dependent contribution i.e. it is too high. Thus, the calculated resistivity value will be lower than the real dc resistivity value. The results of all specific volume resistivity measurements (including the Ic[60 seconds] values) are plotted in Figure 5.6. The curves show that the difference between the I(dc)-data and the Ic(60)-data is mainly present if the sample is in its glassy phase. For temperatures higher than 29oc the time dependent part of the charging current has vanished within 60 seconds i.e. Ic(60) and Idc are equal.

The volume resistivity/temperature curve clearly illustrates the strong difference in the temperature dependency of the volume resistivity of a polymer in its glassy phase and in its rubbery phase. The Tg-value, obtained by drawing two tangents near this glass-rubber transition, is determined at 1000/T = 3.58 or 6"C. This Tg-value is a real static Tg-value and its hypothetical frequency [f(h)] in the frequency/temperature plane will be lower than the f(h) = IxE-2 to ixE-4 claimed by Phillips [12] for dilatometric experiments. A good fit on the Arrhenius plot was obtained assuming an f(h) of ixE-6 for this Tg-value, see below.

The dielectric loss data calculated from the I(discharge) currents according to equation 5.19 show a clear maximum due to the glass-rubber transition for the experiments at 29~ and 20~ These data are plotted in Figure 5.7 as a function of the frequency together with the AC data. The results of both measuring systems can be fitted reasonably. The measured dielectric loss and tangent delta maximum temperatures are

Table 5.1 The dielectric loss/tangent delta maximum temperatures of a liquid epoxy/polyamide resin system

- ' ~ " " = ' " ' . l IT' .... ' , ,~ ,

~ measuring I frequency, Hz. | (Ln ~}

,, ,, ,

AC results 1 MHz. (15.65)

400 kHz. (14.74) i00 kHz (13.35) 40 kHz (12.43)

I resu i | 1 . 0 x e - 3 - 5 7 i 1. ?xB-4 . o

) - 8 4 )

1.0xE-6 1-Ii 98)*

I n ........... ! ,I,

dielectric loss maximum temperatures, oc

, ,,

101.5 95.5 88.5 85 76.5

29 20 6

.... . . . . . . . . . . ~ . . . . . . . . : -- : . _ ' , , 1 " ~ ,,,r . . . . . . i f l

i l

tangent delta | maximum i temperatures, ~

94 87.5 81 75.5 71

,, ,, , , ,, " ! ~ , , , ~ L,J~, . . . . , , , , , . . . . . . . . . , , , , , , ._ ,~ , : .~ , , , , , , = , . . . . . . . , , , , _

* estimated, hypothetical frequency f(h).

138

Figure 5.6 Specific volume resistivity/T(-1) relation of a liquid epoxy/polyamide resin system

+ /(dc) ~ Ic(60) da ta data

l e + 16 .--

l e + 15

E l e + 14 =. ,

c"

0 -

l e + 13 .-=- >:, �9 . , . , - - . . > :~ l e + 12 o , _ . .

"- l e + 1 1 @ E

l e + 10 --- 0 >

._o l e + 0 9 4 - - . . . , . , , -

0 (D EZ

09 1 e + 0 8

l e + 0 7

1 0 0 0 0 0 0

+ /

/ 2 . 5 0

. t . . + ~ - + ~ - " + +4

p*/ z~

/ 4~ /

+ / +

/

. I j

2 . 9 0 , , I J I , I I ,

.3 .30 .3 .70 4 . 1 0 4 . 5 0

1 0 0 0 / T , K ,,, ( - 1 )

139

Figure 5.7 The dielectric loss/frequency relation of a liquid epoxy/polyamide resin system

+ temp. z~ temp. 2 9 ~ 2 0 ~

2 .20

2 .00

1.80

1.60 cb {b o 1.40

0 1 . 2 0

v

0 ~- 1.00 0 r

0 . 8 0 a

0 . 6 0

0.40

0 .20

0 .00

/ +

/ -+

A %

J

A \

\ \ \

\ \

17

15

13 11

9 7

5 3 1

--1

--3 --5 --7

--9

--11 --13

2.60

Liquid epoxy/polyester reein ~ t e m Arrhordus plot dielectrio loss data

+ A C dmla D C dam . . . . ~ , ,

I

2.80 3.00 3.20 3.40 3.60

lO00/T(rna)<.). K(-1)

\

" + + ~

, i i~ . i I , i l u , | i i l l l U , I i I i i l i , i ~ i i l i l i l d * , l l l , l l l i l J l * l i ~ * i i l l l l l i _ .... I I , , l i l i ~ _ | t i l l U l l I I l l l l l l

1 0 - 4 1 0 - 3 1 0 - 2 1 0 -, 1 0 ~ 1 0 ' 1 0 2 1 0 3 1 0 ~ 1 0 5 1 0 6

Frequency, Hz.

140

collected in Table 5.1, the dielectric loss data are plotted in an Arrhenius plot inserted in Figure 5.7. The AC and DC results can be covered by a straight line indicating an apparent activation energy of 260 kJ/mole, according to equation 5.15.

The E'0 and ~'~ values might be estimated from respectively the maximum E'r value measured at the lowest frequency/highest temperature and the ~'r value measured at the highest frequency/lowest temperature. The following experimental values were available:

frequency ~'r at -50~ c'r at 85~ 1 MHz. 3.01 5.07

400 kHz. 3.02 5.44 100 kHz. 3.04 6.08 40 kHz. 3.05 6.52 i0 kHz. 3.07 7.18 4 kHz. 3.08 7.57 1 kHz. 3.11 8.11

400 Hz. 3.12 8.48 i00 Hz. 3.14 9.33 60 Hz. 3.15 9.78

The c'r values measured at -50~ slowly decrease with increasing frequencies. An estimated c'm value of 3.0 is not unreasonable. The c'r values at 85~ still increase too strongly with decreasing frequencies to use these values for a E' estimation The E* value of 260 kJ/mole was substituted 0 �9

therefore together with the area of the ~''r(l MHz.)/I/T curve in equation 5.17, resulting in a (c', - c'w)-value of 9.7. Hence, the estimated E' value is 12 7 i e the dielectric

0 �9 �9 �9

constant of this resin system is estimated to vary between 12.7 and 3.0 maximally.

5.1.6 Time/temperature suDerDosition of dielectric results Knowledge of the Tg-value/frequency relation of SSBR car-tyre rubbers was necessary during the search for correlations between car-tyre properties i.e. wet grip/rolling resistance and the rubbers' chemical structures, see 4.2.3. DMA experiments are, however, limited to about 100 Hz. Dielectric measurements were thought to offer a reasonable alternative. A series of dielectric measurements at frequencies between 10 Hz. and 100 kHz. was started on carbon black loaded rubber samples. Accurate dielectric measurements on these carbon black containing, vulcanised rubber samples proved however very difficult due to the fluctuating, high conduction level of these samples. Therefore, one of the experimental SSBR systems was vulcanised without carbon black to avoid these conduction problems�9 The DSC Tg(onset)-value of this sample without carbon black was equal to that of the carbon black containing system i.e. -35~ The dielectric tangent delta values of this SSBR system was measured at eight different, isothermal temperatures and at nine different frequencies between i0 Hz. and i00 kHz.

141

Table 5.2 The dielectric tangent delta of SSBR (vulcanised /no carbon black) as a function of the temperature and the frequency.

~_ ii l ,Hi l HID Bl 1111] . "iimlll i i i i. l ' ' , _

fre Isothermal measuring temperatures que ncy 66 c 51 c 30 c 18 c 3 c -9 c -17 c -29c

_ , , . . , ,

I0 0.060 0.0130 0.0017 0.0009 0.0015 0.0050 0.0130 0.0110 Hz. 33 0.0240 0.0055 0.0010 0.0008 0.0024 0.0080 0.0170 0.0080 Hz. i00 0.0090 0.0022 0.0008 0.0013 0.0040 0.0120 0.0190 0.0054 Hz. 330 0.0036 0.0012 0.0010 0.0020 0.0065 0.0150 0.0185 0.0039 Hz.

1 0. 0016 0.0009 0.0016 0.0035 0.0100 0.0190 0.0150 0.0028 kHz

3 0.0010 0.0010 0.0028 0.0060 0.0140 0.0190 0.0110 0.0023 kHz I0 0.0011 0.0017 0.0047 0.0095 0.0180 0.0155 0.0080 0.0022

kHz 33 0.0019 0.0030 0.0080 0.0140 0.0190 0.0120 0.0062 0.0023

kHz 100 0.0030 0.0055 0.0128 0.0190 0.0170 0.0092 0.0052 0.0028 kHz

, .[ l ,iiii I,I III l m,l ira!" L ' T 'I Ill I I ,I ' ..... l I I l I L ........... II [I]l [.I

The results of these measurements are listed in Table 5.2 and plotted in Figure 5.8. The tangent delta maximum value due to the glass-rubber transition is clearly visible in the curves measured at 3~ -9~ and -17oc.

These data can be used to show that time and temperature effects are often coupled for relaxation phenomena (the time~temperature superposition principle [13]). Effects due to an temperature increase can thus also be obtained by an increase of the experimental time scale. Hence, the curves in Figure 5.8 were shifted along the frequency-axis while the curve measured at 18~ was chosen as reference temperature. The resulting tangent delta/frequency relation of the SSBR sample at 18~ is shown in Figure 5.9. These results indicate that such a shifting possibility offers data extending over more than eleven decades of frequency while the actual measurements were performed over only four decades.

The used shift factors (Log At) are plotted as a function of the temperature in the graph inserted in Figure 5.10. Log At is described by the so-called WLF-equation [13]"

Log At -- [Cl. (T - Tg)]/[C2 + (T - Tg)] 5.22

C1 = -17.44 and C2 - 51.6 are often mentioned as 'universal' values. The Log At/temperature curve of this rubber system does not fit however with equation 5.22.

1 0 -2 +-,

79

4-J C r [3 c

4-'

142

Figure 5.8 Tan delta/frequency/temperature relation of vulcanised SSBR (no carbon black)

+ A O -I" i �9 + Z~ 66~ 5 1 0 C 30~ 18~ 3~ -9~ -17~ -29~

0

0 . . . . .

i5 1 0 -3

�9 1 , ,,

+

, j + ~ / . - ~+~.,,~-.~+- Q + Q

/.j ,~~ \ / %/ ,,/ / / Jx +,',,~ / ~ / + / , -,OZoj ,o , ,

/ _ I I I I I I l l I I I I I I I I 1~ I I I I I I I I I I I I I I

1 0 ' 10 2 10 3 1 0 " 10 5

F r e q u e n c y , Hz.

143

10 - 2

r0

r @ s

- F J

0 . . . . .

k _

0 q) @

1 0 -3

Figure 5.9 Tan delta/frequency relation of vulcanised SSBR (no carbon black)

+ A 0 "1- & �9 + Zl

66~ 51~ 30~ 18~ 3~ -9~ -17~ -29oc

+

- +

A

+

A

o + @

A +

0 + 1 ~

- , . . I . . . . I . .

- 2 0 2

91)

+t-

& §

0 A

++

r %

e

+ A +

A

A A ~ A

A

re fe rence temperature: 18 ~

i _ . I , I , I .I

4 6 8

I i

10 12

L o g f r e q u e n c y , Hz.

144

Figure 5.10 Ln f against 1/T for tan delta maxima of vulcanised SSBR (no carbon black)

+ d/e lect r /c A U . S . / D M A / D S C d a t a d a t a

1 6 -

1 4 -

1 2 -

1 0 -

8 -

'--' 6 X

E 4 - {--

_..I 2 -

0 -

~ 2 -

- 4 -

- - 6 -

- 8

A

WLF shift factor (log At)/tempemture relation for SSBR

3

2 o= . . j 1

o

- 1

2

3 . .

- 3 0

I

+ dielectric data

- I . \ -I"

\ \ \

§

- l O l O 3 o

Temperature, ~ I _ , . I

4"

i

7 0

, I , I , . . I

3 . 1 0 3 . 3 4 3 . 5 8 3 . 8 2 4 . 0 6 1 0 0 0 / T , K ( - 1 )

4 . 3 0

145

These dielectric measurements provided information about the location of the glass-rubber transition in the frequency/ temperature plane, see Figure 5.10. An ultrasonic measuring system was developed later on during a research program on SSBR rubbers, see 4.2. That offered the possibility to measure carbon black containing samples in the frequency region around 1 MHz. The results, in combination with the low frequency DMA/DSC results, are also plotted in Figure 5.10. The differences between the dielectric results and the low and high frequent dynamic mechanical results can be a measuring technique effect and/or a carbon black effect (present in the dynamic mechanical samples - not in the dielectric samples). These differences illustrate that dielectrically obtained Tg- value information from vulcanised rubber samples without carbon black is less suited for chemical structure/wet grip - rolling resistance correlation studies of (carbon black containing) car-tyre rubber systems than dynamic mechanical data.

5.1.7 The dielectric constant of riuid PU foams Rigid polyurethane foam (PUF) preinsulated pipes for district heating systems need measuring systems to detect leakage in a pipeline network while it is in use. One of such systems measures the absorption of a reflected pulse, transmitted via a copper wire embedded in the PUF insulation. Information about the dielectric constant and loss factor of these foams as a function of the foam density and the MDI index (see 4.1.3) was necessary to calibrate such a system.

Two PUF sample series were prepared based on polyol and isocyanate. The first series of five samples contained between 0 and 30 phr. blowing agent resulting in foam samples with densities increasing from 23 up to 106 kg/m3. The MDI index (see 4.1.3) of these samples was Ii0. The second series of four samples consisted of two samples with a MDI index of 90 and two with a MDI index of 125. From each of these two samples one sample was mixed according to the standard procedure, the other was less good or on purpose mixed. Sample disks of 6 cm. diameter and a thickness of 3.5 mm. were machined from these nine PUF samples and used for the dielectric measurements.

The dielectric response of these PUF samples is weak due to the small amount of material between the measuring electrodes. The dielectric measurements were performed, therefore, with a General Radio 1621 Precision Capacitance Measuring System. This system measures capacities between IxE-6 pF and ixE6 pF with a basic accuracy of • 0.001% (i kHz.) and resistances between IxEl6 Ohm to ixE3 Ohm with a basic accuracy of 0.1% (I kHz.) in a frequency region from 20 Hz. up to I00 kHz.

Accurate sample thickness determination of these thin foam disks is difficult. A guarded measuring cell with two low- potential electrodes (one disk electrode for the sample and one ring electrode for the sample thickness determination) was

146

i sample

~//i11111111111111/11/ (::'2 C1

guarding

temperature measurement

Figure 5.11 Sample cell with capacitive sample thickness determination

147

used, see Figure 5.11, to solve this problem. Cells like this were used earlier to measure the dielectric properties and the linear thermal expansion of polymers simultaneously [14, 15, 16]. The value of capacity C2 in this system provides the sample thickness information after a proper calibration. The C2 capacitance value versus electrode distance relation was calibrated using calibrated spacers and proved to be slightly frequency dependent:

20 Hz. - 1 kHz.: d = 19.9388/(C2 + 0.3435) I0 kHz.: d = 20.0848/(C2 + 0.3866)

i00 kHz.: d = 20.1234/(C2 + 0.3375)

where, d = the electrode distance, ram. C2 = the capacitance, pF.

The weight of the low-potential electrode slightly impressed the thin foam samples. Hence, capacity C2 was measured before and after the sample capacity (Cl) and conductance (GI) measurement. The average sample thickness during the CI/GI determination was calculated using both C2 capacity values. These measurements were performed at 22~ • I~ and a relative humidity of 50%. The results of these measurements are collected in the Tables 5.3 and 5.4.

The dielectric constant values between 20 Hz. and i0 kHz. show a small but clearly present decrease due to the increasing frequency. All the I00 kHz. values are however slightly too high probably due to stray capacitance effects. The dielectric constant/density relation proved to be linear and can be described by-

E'r = 0.00172 x (density) + 0.987 (f = 20 Hz.), E'r = 0.00166 x (density) + 0.988 (f = 1 kHz.), E'r - 0.00165 x (density) + 0.987 (f = i0 kHz.).

The Rval.-values of these linear relations are higher than 0.9999 and the dielectric constant values for a density value of zero fairly approach the theoretical value of 1.000. The linearity of this relation is also reported in literature [17, 18]. Figure 5.12 shows that the four samples with a MDI index of respectively 90 and 125 fit the index Ii0 data. This means that the composition differences (chemically - MDI index or mechanically - bad mixing) do not detectably influence the dielectric constant i.e. the density is the main parameter influencing the dielectric constant of PUF.

The tangent delta values measured at 1 kHz. and i0 kHz. also increase linearly with the density, see the plot inserted in Figure 5.12. The 20 Hz. tangent delta values, however, show a strong non-linear behaviour. A closer look at these data shows that this non-linearity is caused only by the measured value for the sample with the highest foam density. The strong tangent delta increase for this sample is probably caused by an additional DC conduction effect. The tangent delta/ density relations are described by.

0 4~

0

"0

C O0

-n

0

o.

d)

(2_.

E] o~

CO

0

Die

lect

ric

cons

tant

b b

b b

b L~

x

L~

L~

"~

k~

' I"

'

I '

I '

I '

"I

' '

I '

I '

' I

' I

' '

_

.... ~

"\

...

....

....

. "I

.~~

b

o o

C "8

r lh

PO

0 -

0 bO

4~

0

_s

4-

II~

--I -n

I>

= heb

i,,~

0

0 ~-.

149

tan 6 - [4.23 - 0.034x(density) + 0.003x(density)^2] .E-4 (f = 20 Hz.),

tan ~ = [0.18x(density) - 0.43] .E-4 (f = 1 kHz.), tan ~ = [0.23x(density) + 4.13].E-4 (f = I0 kHz.).

The dielectric properties of these foams are not only influenced by the foam density but also by moisture absorption and temperature effects. A single scouting experiment showed, that the increase of the dielectric constant due to heating to 100~ is small in comparison with the density dependency (a 45.7 kg/m3 foam sample was heated from 22~ to 100~ the dielectric constant at I kHz. increased from 1.065 up to 1.071).

Table 5.3

, I ....

density kg/m3

.~..

i 23.0

i 37.1 || ......

45.7 �9 | ,

53.8 ,|

j ~o6. o [

81.2

83.5 ,,

68.3

" 6 7 . 1 . . . . . . . . .

150

The PUF dielectric constant/density relation at 22~ R.H.

MDI index ,,

Ii0

110 ,,

ii0 ,, ,,

ii0

Ii0

125

125 , ,

9O

9O . . . .

,.,

E'r 20 Hz.

,

1.027

1.051

1.066

1.079

1.170

1.122 ,

1.129 ....

1.106

1.105

c'r i kHz.

..... ,

1.026 ,,

i. 049 . . . .

1.065

1.077

1.164

1.120

1.126

1.104

1.101 , ~i

E'r i0 kHz.

1.024

1.048 ,

1.063

1.075

1.161 �9 ..

1.117

1.122

1.101

1.099 P

E'r 100 kHz.

,,

1.028

1.051

1.066

1.077

1.162 ,

1.119

1.125 ,

1.104

1.100 '~' I ....

Table 5.4 The PUF tan ~/density relation at 22~ R.H.

, .., ,,.~ . . ,

density .. ,kg/m3

23.0

37.1

45.7 m

53.8 m

106.0 |L

81.2 ,.,

83.5 ,|

68.3 ..

67.1 ,,.

MDI index

Ii0

110

110

110

ii0

125

125 ,,

90

90 , ,,, I, ,

,,,,

tan 6 20 Hz.

4.7E-4

7.8E-4

9.1E-4

I0.6E-4 , ,

34.6E-4

14.0E-4 �9 ,

20.8E-4 ,

18.3E-4 .

30.8E-4 'rl , , , , , , ,r

tan 1 kHz.

,,

4.1E-4

6.3E-4

7.1E-4 ,

8.9E-4 ,

18.6E-4 , .,

12 .IE-4

13.1E-4

II.9E-4

11.8E-4

tan 10 kHz.

9.2E-4

13.0E-4

15.0E-4

16.9E-4

28.9E-4 , ,

22.2E-4

23 .IE-4 , ,,,,

21.7E-4

21.2E-4 , ........

tan 100 kHz.

12.2E-4

16.8E-4

22.2E-4

24.8E-4

54.5E-4 , ~,,,,

37.3E-4 . . . . .

38.7E-4 ,

37.4E-4

34.6E-4

151

5.2 Effect of moisture on the electrical properties of polymers

5.2.1 Introductio~ Nearly every polymeric system absorbs some moisture under normal atmospheric conditions from the air. This can be a difficult to detect, very small amount as for polyethylene or a few percent as measured for nylons. The sensitivity for moisture increases if a polymer is used in a composite system i.e. as a polymeric matriX with filler particles or fibres dispersed in it. Water absorption can occur then into the interfacial regions of filler/fibre and matrix [19]. Certain polymeric systems, like coatings and cable insulation, are for longer or shorter periods immersed in water during application. After water absorption, the dielectric constant of polymers will increase due to the relative high dielectric constant of water (80). The dielectric losses will also increase while the volume resistivity decreases due to absorbed moisture. Thus, the water sensitivity of a polymer is an important product parameter in connection with the polymer's electrical properties. The mechanical properties of polymers are like the electrical properties influenced by absorption of moisture. The water sensitivity of a polymer is therefore in Chapter 7 indicated as one of the key-parameters of a polymeric system.

The water absorption of composite systems and the effect of this water on the electrical properties was the subject of many studies [2, 20, 21, 22, 23]. Cotinaud, Bonniau and Bunsell [24] observed three mechanisms of water absorption. The first mechanism is the reversible Fickian diffusion of water molecules into the matrix, causing a slight increase of the dielectric constant and the dielectric loss. The second mechanism of redistribution and regrouping of absorbed water molecules is observed at higher humidity levels and causes a strong increase of the dielectric losses and the electrical conduction. The third mechanism, which only occurs on immersion, is transport of water along microcracks in the matrix material through capillary action.

The influence of moisture on the dielectric properties of three experimental resin casting systems and an epoxy based laminate is investigated in this chapter to see if the mechanisms described above, can be recognised. Besides, the resistivity of an epoxy based tank coating and that of plasticised PVC cable insulation material in contact with water is described.

152

5.2.2 Influence of moisture on the dielectric properties of _ _

resin castings and laminates. The effect of moisture uptake on the dielectric constant and the tangent delta is shown for two MDI (4,4'-diphenyl methane diisocyanate) based resin systems, cured with 2% DMP 30 (tris[dimethyl-aminomethyl]phenol), as a function of the temperature in Figure 5.13. Both resin casting samples needed about 60 days to reach their equilibrium water saturation during storage at 20~ and a relative humidity of about 70%. The MDI/lactone resin reached under these conditions a moisture concentration of 1.7 %wt. while 1.3 %wt. was measured for the MDI/styrene system. These 'wet' samples were measured from -120"C up to 240~ at a heating rate of l~ in a nitrogen atmosphere. The moisture absorbed is released during this heating procedure i.e. during the subsequent cooling scan the properties of the dried samples are measured. The increase of the dielectric constant due to absorbed moisture is clear for both systems. These measurements also showed that replacing the lactone by styrene improved i.e. decreased both the dielectric constant and the tangent delta values of this casting system.

In contrast with the dielectric constant, the tangent delta/temperature relation between 0~ and 240"C is not (detectably) influenced by these moisture concentrations. The tangent delta/temperature curves of both 'wet' systems show, however, a clear relaxation effect with maxima between -60~ and -70~ These effects disappear after drying i.e. they stem from the water phase. Such a low-temperature, low-frequency loss process was also detected in dynamic mechanical measurements. Banhegyi et al. reported such an effect due to absorbed water for CaC03 filled polyethylene samples [21].

Wagner showed that conducting particles, dispersed in a non- conducting matrix can cause an energy dissipation maximum, the so-called Maxwell-Wagner dielectric relaxation. Absorbed water in a polymer might considered to behave like conducting 'particles' in the non-conducting polymer matrix. We tried to apply, therefore, this theory to calculate the dielectric properties of these 'wet, resin systems. The relaxation time of such a Maxwell-Wagner absorption process is according to Hill [6] given by:

T = (2.E'r + E'W)/(4.~.90EI0.T W) 5.23

where: E'r = MDI/lactone E'r at 20~ and ikHz., dry sample i.e. 3.39

E'w - dielectric constant of water (80), T w ~ the water conductivity, (Ohm.cm) ̂ (-i), T = relaxation time, s.

The use of equation 5.23 is hampered by a lack of knowledge about the conductivity level of the absorbed moisture phase. Equation 5.23 was therefore used to calculate this absorbed water conductivity after an estimation of relaxation time T. The experimental activation energy value of the dielectrically

153

5.0

4.0

3.0

MEASURING FREQUENCY IS IOOO HERZ

SECOND SCAN, DRIED SAMPLES

2 . 0 -

v MDZlLactone 2/1 a MDZlStyrene 2/1

I i 1 ~ i I ,,,a~ I I I . I , I �9 I . I , , , , I

-IZO -80 -40 0 40 80

DIELECTRIC CONSTANT - 5.0

- 4.0

- 5.0

-ZO

, 1.0 120 160 2 0 0 2 4 0 28o

TEMPERATURE (CENTIGRADES)

INFLUENCE OF MOISTURE UPTAKE DURING STORAGE UNDER ATMOSPHERIC CONDITIONS ON THE DIELECTRIC CONSTANT

IO

.TANGENT DIELTA

r MEASURING FREOUENCY IS IOO0 HERZ

L w

DRIED SAMPLE, SECOND SCAN

V MDI/Lactone 2/1 A MDIIStyrene 2/1

-120 - 8 0 - 4 0 0 4 0 8 0 120

i0 -I '2

w m

f -

_ I f f "s =, , i m

m

I l y 4

t60 200 240 200 TEMPERATURE (CENTIG RADES)

INFLUENCE OF MOISTURE UPTAKE DURING STORAGE UNDER ATMOSPHERIC CONDITIONS ON THE TANGENT DELTA

Figure 5.13

154

and dynamic mechanically measured moisture loss maxima, plotted in an Arrhenius plot, was 46 kJ/mol. (Illers [25] reported a value of 54 kJ/mol for the water relaxation effect in nylons). An activation energy of 46 kJ/mol means that a 'moisture' relaxation maximum, measured at -60~ using a frequency of IkHz. will shift to 20"C if the measuring frequency is increased to 1.22 MHz. Hence, relaxation time T might be I/1.22E6 = 8.2E-7 s at 20~ Substitution of this value in equation 5.25 results in a value for the water conductivity of:

T (absorbed moisture) = 9.4E- 6 (Ohm. cm) ̂ (- i) �9

This seems to be a reasonable value. Hill [6] is using a conductivity value of 1.0E-6 (Ohm.cm)^(-l) for pure water. Steeman et al. [19] are using a value of 5.0E-8 (Ohm.cm)^(-l) for the bulk conductivity of pure water at 20~ It is questionable, however, if absorbed moisture can be considered as pure water. Ionic impurities from the resin matrix can easily increase the conductivity of the water phase.

The dielectric constant of the 'wet' resin is according to the Maxwell-Wagner theory:

~'(U) = c'r[l + 3.U(E'W- E'r)/(2.E'r + E'w)] 5.24

a n d

E' = E' (U) [i + S/(I + ~'.T2)] 5.25

and

S = (9.~.E'r)/(2.E'r + E'W) 5.26

where- u = the volume fraction of moisture i.e. 0.02 and ~' = the dielectric constant of the 'wet' resin i.e.

3.84.

Equation 5.24 gives an c' (u) value of 3.57, equation 5.25 reduces for the given w and T value to:

E' = E' (~). [i § S] = 3.57. [i § 0.007] = 3.60

The dielectric constant at 20~ increased from 3.39 to 3.84 due to the 1.7 %wt. moisture (2.0 %v.). The calculated increase of the dielectric constant from 3.39 to 3.60 is only about 50 % of the total effect. The Maxwell-Wagner theory thus seems to describe roughly the frequency/temperature location of the dielectric loss maximum due to absorbed moisture. However, it does not adequately describe the increase of the dielectric constant due to the moisture uptake from the air. A possible reason for this discrepancy might be that one of the assumptions does not hold, viz. that the conductivity of the resin matrix is negligibly small.

155

Brasher [26] adapted the Hartshorn equation to describe the increase of the dielectric constant of coatings due to water absorption-

~' = E'r. (~'w.)^(n.u) 5.27

where- u = the volume fraction of water and n -- an adjustable parameter, which characterises

the effectiveness of a certain water concentration i.e.

n > i, water absorbed in pores parallel to the direction of the electrical field applied,

n = 1, water-filled spherical interstices, randomly distributed in the sample and

n < 1, water bound to the resin:

Equation 5.27 describes the increase of the dielectric constant of the MDI/lactone system from 3.39 to 3.84 due to 2.0 %v. of moisture using a value of 1.42 for n. The dielectric constant of the MDI/styrene system at 20~ increased from 2.94 to 3.24 due to 1.56 %v. of moisture. This increase also corresponds to a value n = 1.42 in equation 5.27.

A DGEBA/HHPH (diglycidyl ether of bisphenol A/hexahydrophthalic anhydride) casting system reached an equilibrium moisture concentration of 0.56 %v during storage at 20"C and 70 % R.H. The dielectric constant increased from 3.31 to 3.51 due to this absorption effect. This increase is corresponding to a value n = 2.39 in equation 5.27. The sample was subsequently immersed in demineralised water until a new equilibrium saturation of 1.08 %v. was reached. The dielectric constant increased from 3.51 to 3.85 due to this absorbed water additionally. The results of this experiment are plotted in Figure 5.14. The drawn line corresponds to equation 5.27, using the n value of 2.39. The calculated and measured dielectric constant values are equal up to a water content of 0.71%v. The measured dielectric constant values increase, however, stronger than those calculated at higher water contents. The same effect is measured for the tangent delta values. This gives the impression that we observe the mentioned transition in the introduction from the second to the third mechanism. The transition from the first to the second mechanism is not detectable due to a lack of data points at low moisture contents.

Figure 5.15 shows similar results for an epoxy based glass fibre laminate (Nelco-100, a one-side copper-clad laminate). The moisture content scale in this figure is an experimental one i.e. it is the weight increase due to moisture absorption divided by the~~11weight of resin, glass and copper. Although the moisture content scales of both figures 5.14 and 5.15 can not be compared, the change in the slopes of both

156

Figure 5.14 Influence of moisture on the dielectric properties on a D G E B A / H H P A casting

+ die/. A ta ngen t c o n s t a n t de / ta

~q 3: x/

T " -

"0 c -

0 0

0 0,1

c 0 O

0 ~

0 (D r

4 . 0 0

3 . 9 0

3 . 8 0

3 .70

3 .60

3 .50

3 .40

: 3 . 3 0

3.20

3 . 1 0 z~-

: 3 . 0 0

calculated curve Hartshorn equation

0 . 0 0

J +

I ,

+

j Z~ /

/ / A

+

/ /

/

/

A / /

/ /

I , i , , I J I i I

0 . 2 0 0 . 4 0 0 . 6 0 0 . 8 0

Z~

0.1

- T

O

o " 0q

~3

-(3

c �9

c

, 0 . 0 1

1.00

Mois ture content , %wt .

157

Figure 5.15 Influence of moisture on the dielectric properties of an epoxy based laminate

+ die/. z~ t a n g e n t c o n s tan t d e l t a

6 . 0 0

3 : 5 . 8 0

" T "

-(9 E 5 .60

0 0

0 5 .40 Od

5 .20 E

r E 0 o 5 .00 0

0 --r 4 .80 @ . . , ~ .

db

4 .60

J

L~ ~

I

0.00

/+ +

+

/

+7+- j z~

, I , i , . , I I ,,I

0 .20 0 . 4 0 0 . 6 0

I

73 C

- 0 . 1 0 0

. 0 O~

-

- r -0

. E O~ E

F- 0.01

0 . 0 0 5 0 . 8 0

Mois ture content , %wt .

158

resin casting curves is also clearly present in both curves of the laminate sample. Especially the inflection in the laminate tangent delta/moisture content curve is more pronounced than that in the resin casting curve. This might be caused by the relative high water/resin ratio in the laminate sample (the overall moisture concentration of 0.5 %wt. in Figure 5.15 - the inflection concentration - is estimated to be about 1.3 %wt. on pure resin). Alternatively, the distribution of water absorbed water in a glass fibre laminate sample might differ from that of a resin casting system. Woo and Piggot [27] state, for instance, that absorbed water in a glass-epoxy composite is concentrated in the interfacial regions, which are interconnected by disk-shaped water clusters providing conducting paths.

5.2.3 The effect of seawater and crude oil on the electrical properties of a tankcoating system~ The vast increase in the bulk shipment of refined and chemical products has resulted in a growing interest in coatings for tanks having an improved resistance to various chemicals. The electrical properties of such systems are important in connection with: - the corrosion protection, - the decay of electrostatic charge and, - as an aid to monitor changes in the coating's performance.

The principal aim of actions taken with regard to the electrostatic problems has been to promote the safe discharge of such charges during their development. The two factors which determine the ability of electrostatic charges to disappear through the paint layer are its volume resistivity and dielectric constant. The product of these, called the 'relaxation time', represents the time taken for the charge to decrease to i/e th ('e' the base of the natural logarithm) of its original value, and is used as a measure for the coating's charge dissipation ability. If the resistivity and the dielectric constant of a tank coating are of the same order of magnitude as those electrical properties of the product carried, the relaxation time of the whole system is not increased and this situation is considered acceptable.

From an anti-corrosion point of view, however, coatings with a high volume resistivity are favoured. The resulting high resistance of the paint layer can act as a barrier preventing ions reaching the metal surface; this is called resistance inhibition. This factor will result in a greater disparity between the electrical properties of the coating and typical cargoes and thus tends to increase the likelihood of electrostatic charge built-up. Apart from this aspect, the volume resistivity may also considerably affect the efficiency of the (impressed) current cathodic protection systems typically used in tanker situations.

159

A further complication is that the volume resistivity of coating materials is strongly dependent on their composition and this, in turn, can be greatly affected by contact with different cargoes. Hence, knowledge of this property as a function of contact time with different media is desirable. The electrical properties of a series of eight different tank coating systems in contact with respectively seawater, crude oil and kerosine were determined [2], as a part of a tank coatings' research program. The results measured for one of these coatings, an epoxy coal-tar system cured with an amine adduct (26 %wt. binder, 35 %wt. coal-tar and 39 %wt. talc/ barytes pigments) are reported in this chapter.

Sample cells (Figure 5.16) were constructed upon the coated panels by glueing the polished end of glass tubes (internal diameter about 50 mm.) onto the coated surface with epoxy cement. The sample cells were stored at 22 • 2~ and a relative humidity of 55 • 5 per cent. The steel panel served as high potential electrode (H), and a mercury electrode, connected by a platinum wire, as low potential electrode (L). Guarding was achieved by using a bell-shaped brass cover which also supports the connector of the low potential electrode.

These sample cells were filled with the appropriate liquid during the immersion experiments. Before each measurement the liquid was decanted and the surface to be measured was wiped dry with a soft tissue. The mercury electrode was then immediately introduced. After the measurement the mercury was removed and fresh immersion liquid was reintroduced.

Figure 5.17 shows the decrease of the volume resistivity and the increase of the dielectric constant as a function of the square root of the immersion time, due to contact with seawater. The volume resistivity decreases in about six days from 4.2E14 Ohm.cm to 3.3EII Ohm.cm and stabilised finally at a level of 2.2EII Ohm.cm. The dielectric constant increased from 5.49 to 9.36 in about six days and stabilised at a value of 10.25.

The electrical properties of this epoxy coal-tar system hardly changed due to contact with Middle East (Kuwait) crude oil:

vol. resistivity, diel. constant, 0hm.cm i kHz.

epoxy coal-tar system a. before immersion 7.8E13 5.28 b. after 1300 hours of

contact with crude oil 7.2E13 5.30

c. crude oil as such 2.1E 9 2.47

These results and the results in Figure 5.17 show that seawater penetrates into the coating and influences, subsequently, the electrical properties strongly. The crude oil, however, does hardly affect the coating properties.

160

PL AT INUM CONTACT

BRASS GUARDING-

LIOUID MERCURY ELE

GLASS-TUBE, ELECTRODE SURFACE ABOUT 20 cm = \

PA INT COAT INS

H I DC SUPPLY

TO ELECTRO- METER

L ---.~

Figure 5.16 Construction of the sample cell

161

Figure 5.17 Epoxy/coal-tar coating, in contact with sea-water

+ volume A d ie lect r ic res is t iv i ty constant

~

l e + 1 4 0

J �9

~ l e + 1 3

�9 / o 1 e + 1 2 > .

+

X + ~ ~ +_____.__+ ~ ~ + ~

d

u9

l e + 1 1 1 , , , I , . . I , . . . . I i I , I . i I , I ,

11

9 .o 0 O4

- 8

7

6

0 5 10 15 2 0 25 30 35 4 0 45 5

c

c 0 s (J

4-.., (3 G) (b d3

Seawate r immersion time, h~ Q.5

162

Figure 5.18 Epoxy/coal-tar coating, in contact with sea-water followed by crude oil + s e a - A c r u d e

w a t e r o i l

l e + 1 4

E 0

0 >~ >

l e + 1 3

r E D 0 >

0 q - -

O r l e + 1 2 o U3

7-1-

- +

i

2 e + 1 1 , , ,

- 1 8 - 1 4 - 1 0

+

I , L

- 6

S e a w a t e r / c r u d e

i I

A /

A

/

S / A

. . t , I , . I , I ,

- 2 2 6 10 14

oi l i m m e r s i o n t ime , I-1~0.5

163

A nearly complete recovery of the coating's volume resistivity is measured, if an epoxy coal-tar coating after about fourteen days of contact with seawater is immersed in crude oil (Figure 5.18). This indicates that the seawater absorbed is nearly completely leached out by the crude oil i.e. crude oil absorps seawater considerably stronger than the epoxy coal-tar coating does.

The relaxation time T, in connection with electrostatic charge /discharge processes is defined as:

T = c'o.c'r.G 5.28

where: T = relaxation time, s 'o = 8.85E-12, F/m 'r = die1. constant coating,

= volume resistivity coating, Ohm.m.

Equation 5.28 results in a relaxation time of about 50 seconds for the epoxy coal-tar coating as such. This value decreases to about 0.2 seconds after contact of the coating with seawater. The relaxation time of the epoxy coal-tar system remains higher than the relaxation times of 0.018 s. to 0.00018 reported for different crude oils [28], even after prolonged contact with seawater.

5.2.4 The Ki-value determination of PVC cRble compounds The specific volume resistivity is the most important electrical property of an electrical grade PVC. It is measured on a heavily plasticised product, the 'cable compound', pressed to a 2 mm. thick sample sheet. Cable manufactures usually test the resistivity of these compounds on cable samples and express the results in a so-called Ki-value. The Ki-value is in fact a volume resistivity value (see below) but measured on a cable sample with tapwater as low potential measuring electrode. A series of Ki-value determinations was performed to investigate the different parameters influencing this quantity.

The Ki-value determinations were performed on coils with a diameter of about 0.2 m., made of cable samples with a length of about 10 m. Such a coil is immersed in an (electrically insulated) water-bath filled with tapwater and kept at the specified temperature. The electrical resistance is then measured between the copper wire of the cable (high potential electrode) and the tapwater (low potential electrode). The Ki- value is calculated according to [29]-

Ki - Ri/[log(R2/Rl) ] 5.29

where- Ki - insulation constant, Mega-Ohm.km Ri = insulation resistance, Mega-Ohm.km R2 = cable diameter, mm R1 = wire diameter, mm.

"1"1

m

m.

CQ

C

CD

~D

ii

\

!

165

We measured the voltage over, and the current flowing through a piece of cable of about I0 m. instead of one kilometer length. Hence, the Ri-value follows from-

Ri = (V.I.E-6)/(I.E3) -- (V.I.E-9)/I 5.30

where : Ri = insulation resistance, Mega-Ohm.km V = measuring voltage, Volt I = measured current, Ampere 1 = sample length, m.

Substitution of 5.30 in 5.29 is giving the Ki-value according to:

Ki = [V.I.E-9] / [I.Iog(R2/RI) ] 5.31

The (in this way) calculated Ki-value in Mega-Ohm.km has in fact the same dimensions as the specific volume resistivity which is expressed in Ohm.m. The derivation of Ki starting from equation 5.2 is straight forward: the resistance of a rectangular piece of material between two flat metal electrodes is given according to equation 5.2 by:

R = (G.L)/A 5.32

where: r = specific volume resistivity, Ohm.m R = resistance, Ohm L = length in the direction of the current, m A = area perpendicular to the current direction, m2.

In the case of a cable, however, an annulus at a distance r from the center of the cable and having an infinitesmal thickness dr, is considered (see Figure 5.19). The length of this annulus in the direction of the current is dr. Its cross section perpendicular to the direction of the current is 2.~.r.1, where 1 is the length of the cable. The resistance dR of this annulus is then according to 5.32:

dR = (o.dr)/(2.~.r.l) 5.33

The total resistance of the cable becomes:

R2 R = a/(2.~.l).;dr/r = (G.Ln[R2/RI])/(2.~.I)

R1 5.34

This resistance value is converted into a Ri-value in Mega- Ohm.kin by substitution of 1 = 1000 m and expressing of the Ohms in Mega-Ohms i.e. xE-6:

Ri = (2,30.E-9.a.Log[R2/RI])/(2.~) 5.35

Substitution of 5.35 in equation 5.29 gives for Ki"

Ki = Ri/(Log[R2/RI]) = 0,366.E-9.G 5.36

166

Figure 5.20 Ki-value/temperature relation for a cable sample ex-Dorlyl

E

c- O

I E~ (D

(U > I

IOQ

1 0 ;

-1 .-

0.5 15

Ki VALUE, MO. in 4 5 0 -

410

3 7 0

330, VOLTAg( (I.F.r 10N TIM(

�9 I I I I I , I 40 SO 120 160 200 240 200

IMMERSION T I M ( , h

I r r IT~ rs i on t ime : 2 h o u r s E l e c t r i f i c a t i o n t i m e : 2 rain.

V o l t a g e : 5 0 0 V

I I I , 1

35 55 75 I

95 115

Temperature, deg. C

167

where- Ki = insulation constant, Mega-Ohm.km = specific volume resistivity, Ohm.m.

The Ki-value and G are thus in theory linearly related. The use of a water electrode (resulting in water penetration and plasticiser dissolving effects) makes this relation, however, more complicated.

The main variable during these type of measurements is the temperature. Figure 5.20 shows that the Ki-values of an ex- Dorlyl cable, decrease from 331 Mega-Ohm.km at 20~ to 1.2 Mega-Ohm.km at 80~ It is hence not surprising that the measuring temperature of the Ki-value is specified by every cable manufacturer.

A second measuring variable is the immersion time. The inserted graph in Figure 5.20 shows the effect of the immersion time on the Ki-value at a temperature of 20~ The Ki-value immediately after immersion (428 Mega-Ohm.km) decreases to 330 Mega-Ohm.kmafter two hours of immersion and to 308 Mega-Ohm.km after five hours of immersion. Then the Ki- value starts to increase again to 420 Mega-Ohm.km after 220 hours and to 570 Mega-Ohm.km after 850 hours when this experiment was stopped. The Ki-value decrease might be a moistening effect. The subsequent 'recovery' of the resistance is thought to be caused by depletion of charge carriers due to the water penetration and/or perhaps some extraction of plasticiser. This immersion time effect might be one of the reasons for the relative large differences in the Ki-value results from different suppliers (the immersion time is hardly specified).

In chapter 5.1.2 was shown that the electrification time affects the resistivity determination of polymeric materials. Every manufacturer is using its own (again seldom specified) electrification time for the Ki-value determination. In this case, the situation becomes more complicated: the electrification time effect proves to be immersion time dependent (figure 5.21).

The Ki-values of the Dorlyl cable sample were measured at 20~ using different electrification voltages between I00 V and I000 V. These measurements showed that this variable has no significant influence on the Ki-values at 20~ Some electrification voltage effect might be expected however at higher temperatures.

Table 5.5 lists the results of the Ki-value determinations on three ex-Dorlyl cable samples. These values show that a repeatability of + I0 % is possible under these conditions.

Three different cable samples were, subsequently, subjected to a series of Ki-value determinations at 60~ The results of these measurements are plotted in Figure 5.22. In spite of the scatter, especially in the Silec results, the difference in electrical performance between the three cable samples is

168

Ki VALUE, M~.km 650 -

570 fcAeL~ Ex OORL','L ' / " IVOLT AC3E: eX)O v . f IT,.PeRATu~: =0% / "

490 II 2 h I J =sh . / | _ e = 1 8 h , ,

410

330

, J ,, ,I . . . . I z~-,o~ ,oo ,ooo ,ocx)o

ELECTRIFICATION TIME, $

Figure 5.21 The Ki value of a PVC cable compound as a function of the electrification time

169

120

I 0 0

80

60

40

20

K i ~O, LUE, M,0.. km

B |

IT(MP(I~A"I;:( : eo o(E'-

/ " i ELECT~'C, T=E= ....... 2 ~. SILEC CABLE A

II & A ~

A

A

DORLYL CABLE o

. . . . I . . . . . . . . . . . . . . . . I . . . . . . . . . I I0 tO0 tOO0

IMMERSION TIME, h

Figure 5.22 Ki-values vs immersion time at 60~

170

clear. A Ki-value of 65 Mega-Ohm.km was measured for the Silec ~ sample after two hours of immersion and an electrification time of two minutes; Silec reported a value of 52 Mega-Ohm.km. A Ki-value of 40 Mega-Ohm.km was measured for the CGFCE sample under the same conditions; CGFCE reported a value of 48 Mega- Ohm.km. No additional information about the measuring conditions was available for both samples. The agreement between the measured Ki-values and the by Silec and CGFCE reported values seems to be not to bad considering that both manufacturers only specified their measuring temperatures.

The above described results indicate clearly the importance of a proper specification of the most important Ki-value measuring conditions i.e. - the temperature, - the immersion time, - the electrification time, - the electrification voltage and - the for the determination used cable length. A considerable difference in length (for example 250 m. instead of the i0 m. samples described here) influences the wettability of the sample coil and increases the chance on weak spots in the insulation layer due to the cable extrusion process.

Table 5.5 Repeatability of Ki-value determinations.

Immersion time, h

0

2

5

50 , ,

a. cable sample b. temperature

Ki-value, M. Ohm. km sample .1

428 ,

331

308

380 ,

Ki-value, M. Ohm. km sample 2

398

Ki-value, M. Ohm. km sample 3

466

- 399

- 319

359

: ex-Dorlyl : 20 ~

c. electrification voltage. 500 Volt d. electrification time �9 2 minutes e. sample lenght �9 i0 m.

171

5.3 Conduction improvement of epoxy resins by carbon black addition

5.3.1 Electrostatic safety criteria Polymers are used in many applications especially for their ease of moulding in combination with a high electrical insulation resistance. There are also applications where a too high electrical resistance hampers the discharge of electrostatic charges which can build-up. It is often necessary then to decrease the polymer's original resistivity to lower values for safety reasons. Addition of conducting particles, for example carbon black, is one of the possibilities to decrease the resistivity of a polymer to an acceptable value. This acceptable value is set by rather simple rules:

The discharge process of electrostatic charges is decribed by relaxation time T. This value stems from the charge/discharge processes of a capacitor C via a resistance R, given by.

charging : Uc = Uo. (I - e ̂ [-t/T]) 5.37

discharging: Uc = Uo.e ̂ [-t/T] 5.38

where- T = R.C

For the discharge of a non-conducting object, T follows from:

R = (G.d)/A, see 5.2

C = (~o.c'r.A)/d, see 5.4

T = RC = G.~o.~'r, see 5.28

The resistance R is formed by the insulation resistance to earth of an object, while capacitance C is the capacitance of a standing person i.e. about 100 pF [30]. A T-value of 0.01 s. or less is required for protection of such a person against discharges [30, 31]. The maximum value of R has to be in this case :

R(maximum) - 0.01/I.0E-10 = 1.0E8 Ohm

Distance (d) and area (A) are often difficult to determine. DIN 51953 gives, therefore, as a practical rule: the resistance to earth per 20 cm2 has to be lower than 1.0E8 Ohm. A too low resistance, however, can also be dangerous for persons in connection with the possibility of shortcircuiting. Hence, a resistance to earth per 20 cm2 between 1.0E5/1.0E6 Ohm and 1.0E8 Ohm is usually considered to be acceptable.

172

5.3.2 DC properties of experimental epoxy resin/carbon black systems Polymers are often filled with carbon black to obtain anti- static compounds with specific volume resitivities of about 1.0E80hm.m or less. The resistivity will in general hardly decrease during the first I0 %wt. of carbon black added. Resistivity values of about I00 Ohm.m are realised for carbon black concentrations higher than about 20 %wt., mainly due to direct contact between the conducting particles. The rather small concentration region where the resistivity decreases from a high to low resistivity is called the percolation threshold [32]. However, carbon black loadings of about 20 %wt. seriously reduce the mechanical properties (in a negative sence) .

Some types of carbon black (Ketjenblack EC-2000 and Gulf AB 550-P) in certain cured epoxy resin systems showed to have a percolation threshold of less than 0.5 %wt., see Figure 5.23. The commercial relevance of this knowledge was recognised [33] and additional experiments were performed to obtain more insight in this phenomenon. Some experimental results of this investigation are collected in Table 5.6.

Table 5.6 Electrical properties of carbon black filled epoxy resins cured with different curing agents at 23~

DGEBA resin cured with-

aromatic amine

alifatic amine

cyclo- alifatic amine

DGEBA cure

II

carbon black, %wt.

0.0 1.0

0.0 1.3

0.0 0.9

volume resistivity,

Ohm. m .

2.0El3 2.0E 4

3.1E13 6.2E 7

5.8E12 2.5E12

dielectric constant at 1 kHz.

4.81 45.7 , . .

4.60 25.4

. . . .

4.69 8.38

',,

�9 diglycidyl ether of bisphenol A, : 7 days at 20~

post-cure- 24 hours/100~

Tg-value, DSC onset,

oc

72 66

.,

107 112

. . ..

68 69

These values illustrate the strong influence of the type of curing agent used. The phenomenon studied is clearly present in the first system where the low volume resistivity is accompanied by a high dielectric constant value. This phenomenon is clearly not present in the third, with a cyclo- alifatic amine cured system. The dielectric constant of this system, although increased due to the carbon black addition, is low compared with that of the former system. The resistivity decrease effect is found to some extent present

173

101~

1014

1 0 13

E 12 g l O ..c 0 1011

>:, 10 '0 >

1 0 9 (/) . , , . -

t/') 8 �9 1Q & . .

107 E

10 6 0

105 0

"~ 104

~ 3 ~ 10

10 2

Figure 5.23 Volume resistivity of epoxy casting (DGEBNDDM) versus the carbon black content

!

O

+

10"

10 =

10'

10'

/ ,4- I +

-I-

+

O.O0 0.40 0,80 1,20 1.60

C a d ~ n black c ~ r e l i o r ~ %wI.

2 .00

lO' f 10 ~

0 .00 I . I , , I . I , I I i I .....

O. 10 0 . 2 0 0 . 3 0 0 . 4 0 0 . 5 0

I

0 . 6 0

Carbon black concentrat ion, %wt.

11 10

:z

c)

cK

o zz) ._I 0 >

u . I'-4

LL

U lIJ O. tr~

8 10

7 lO

8

IZ I

5 lO

4

1o ,w-

174

/ / I

f I"

DOEBA/cyclo-alifatic amine cured system (x 0.9 %wt. carbon black) (A no carbon black)

,/ /

x

x

/ X

A

/ /<

/ A

/

X ------- m .....

X ~A /

x~)i A

A / x

DOEBA/alifatic amine cured system (1.3 % w t . carbon black)

Figure 5.24

\ \ +\

+,+

1000/T. KT-1 l l ' ' , , I l

Lr) ~ r~

DOEBA/aromatic amine cured system

~ + . . (1.0 %wL carbon black)

+ ------ § ~ + _ . §

' " t I I I' ' ' I I : I ' - | I " I I .

175

for the system cured with a alifatic amine- the volume resistivity is about six decades decreased and the dielectric constant is significantly increased.

The specific volume resistivity of these three systems was measured as a function of the temperature between 20~ and II0/140~ The results are shown in Figure 5.24. The DGEBA/cyclo-alifatic amine cured system shows a typical polymeric behaviour i.e. the resistivity decreases strongly with an increasing temperature due to an increase of the charge carriers mobility. The DGEBA/aromatic amine cured system behaves like a metal i.e. the resistivity slightly increases with increasing temperatures. In this case this might be caused by an increase of the distance between the conducting carbon black particles due to the difference in expansion coefficient between the carbon black and the resin matrix. This expansion coefficient difference (and thus the resistivity increase with the temperature) increases if the resin matrix changes from its glassy into its rubbery state. The volume resistivity of the DGEBA/alifatc amine cured system exhibits an intermediate behaviour. The volume resistivity increased between 20~ and 100"C as a function of the temperature (metallic behaviour). The increased charge carrier mobility seems to dominate the thermal expansion difference effect, at temperatures above the glass-rubber transition, resulting in a change in the slope of the resistivity/ temperature curve.

Figure 5.25 shows the AC properties (dielectric constant/ dielectric loss factor) at 23~ as a function of the frequency. The differences found in the volume resistivity of these systems are also reflected in the AC properties. The strong decrease of the dielectric constant as a function of the frequency for the system cured with EPIKURE 160/161 indicates a certain amount of capacitive coupling between the conducting carbon black particles. The dielectric loss factor/frequency relation of this system is nearly linear with a slope of about -i, pointing at a pure resistive behaviour (see equation 5.18).

Kawamoto [34] represented a single carbon black particle - polymer interface - carbon black particle 'unit' by a resistor R in series with a resistor R and a capacitor C in parallel. He subsequently assumed, that the whole body of the carbon black filled system can be represented by one single RC circuit. He was able, using this model, to calculate accurately the AC properties as a function of the frequency for a carbon black filled PVC system. This model failed, however, in the calculation of the AC properties of the above decribed, conductive epoxy resin systems.

A variety of microscopical techniques revealed that around some critical concentration well below 2 %wt. of carbon black there is a transition from isolated inclusions to a conducting network [35]. The presence of a conducting network in contrast with Kawamoto's randomly dispersed carbon black particles

" ' " ' " " ' F i g u r e

7a. e + ~ DGEB,~aromatic amine cured system es. �9 X ( I .0 %wt. carbon black)

6g. B §

55. g

58 . g

Temperature = 23~ 45. S

DGEBA/alifatic amine \ cured s3~tem .... (1.3 %wt. carbon black) \

4 8 . I I

35. g

3 8 . g

25. e

2 0 . 8

15. g

l g . g

5 . 8

I N

5.25

DGEBA/cycio-alifatic amine ~ + �9 \ 4,

(0.9 %wt. carbon black) ,,,,, �9

X ~ X ' x X~

M #1 ,e i �9 I I I I "I o,4 l,,e l.,e g,i g,I

4 I l l r

I ~ u u_

0

u

p. I U

t g . ~ w D

8 , l g

8 fg

g l g

-! 1 8

I

"X 4.

+

X 4

DGEBA/aromatic amine cured system (!.0 %wt. carbon black)

Temperature = 23~

DGEBA/alifatic amine cured s y s t e m ~ (I.3 %~. carbon black) 4

% \ "\ \ , " \

~ ~ � 9 ..----A ~ �9 " " A ~ - A ~ §

x - - - - - X -=---X ~ x - ~ / x " x " x 'x X x ~ -

X ~ x ~ X ~ . . ~ X ~ x f x

DGEBA/cyclo-alifatic amine cured system (0.9 %wt. carbon black)

�9 ." : ; ; ":;4; ; ." : ; . '::;; ; : t~HtI~-+~444H----P--F-~4~N ~ ,q im I

I ~ I I m

F R E ~ , H ~ FREQUENCY, H=.

177

may well be the reason that his model failed to describe the AC properties of the conductive epoxy resin systems.

The formation of such a conducting network was found to depend on the surface tension of the used curing agent at the cure temperature. In addition, the degree of dispersion of the carbon black particles proved to be critical [33]. A homogeneous dispersion of the carbon black particles (substantially) having a diameter predominantly below 1.0E-6 m. and a curing agent with a surface tension at the curing temperature of either at least 33 or at most 23 MN/m proved to be necessary for the formation of a conducting network.

It is believed that the aforesaid surface tension criteria are related to the interracial energy between curing agent molecules and carbon black particles. This energy is exceeding a defined critical value, thus creating a driving force which causes the carbon black particles to form a network structure of coagulated particles in a resin matrix.

~.3.3 The DC properties Of anti-static epoxy_ GFR pipes The use of glass-fibre reinforced (GFR) epoxy resin pipes is, especially in tankers, hampered by the 'bad' electrostatic properties of these pipes. The possibility to decrease the volume resistivity to an acceptable level using only a small amount of carbon black (see 5.3.2) resulted in the development of the WAVIMAR anti-static GFR pipe system by Wavin BV. This pipe system is based on a liquid DGEBA/MDA (100/27) and cured for two hours at 120~ containing about 1.5 %wt. (on the resin phase) of Ketjen black EC-2000 carbon black. The specific volume resistivity of a sample of such a pipe was measured as a function of the direction, the field strength and the pipe wall thickness.

The 5.1 mm. thick pipe wall consists of an about 0.5 mm. thick, with C-glass reinforced, liner layer on the inner-side followed by the 4.3 mm. thick, cross-plied glass fibre reinforced, pipe wall body. The outside of the pipe wall consists of an about 0.3 mm. thick epoxy resin layer. The investigated pipe sample had an outside diameter of 50 mm.

First, the specific volume resistivity of a reference sample without carbon black was measured in the radial direction. Figure 5.26 shows the strong time-dependency of the charging current at a measuring voltage of 1200 Volt and a temperature of 23~ The Ic has not yet reached the I(dc) level after even 2.5 hours. The measuring procedure described in 5.1.2 is working satisfactorally and gives a constant I(dc)-value of 2.7E-12 Ampere between 900 and 9000 seconds measuring time. This results in a specific volume resistivity value of 3.0E14 Ohm.m for the reference system without carbon black (measured in the radial direction).

Subsequently, about eighty millimeter long sample pieces (properly mounted and supplied with vacuum evaporated silver

178

Figure 5.26 Specific volume resistivity determination of an epoxy resin based, GFR pipe + I c - z~ I d - o I ( d c ) -

c u r r e n t cu r ren t c u r r e n t

1 0 -11 (1) !1,,,_ (1) o E ,<

E L .

D (3

1 0 -12

10- '~ I::

1 0 -13

10"

~ [ '

\

X:

~ 0 ~ 0 ~ 0

o/~

~ + ~ + _

o - 0

Pipe wall th ickness 5 mm. Measuring voltage 1200 V

0 I . I I I . I 1 I I I I I

1 0 3

I I I i

1 0 "

Charge/d ischarge time, s.

179

electrodes), were used to measure the volume resistivity of the anti-static pipe in both the radial and the axial direction as a function of the electrical field strength. These measurements were performed on the pipe samples as such, on the pipe samples after machining away the 0.3 mm. thick outside resin layer and finally after machining away a layer of 2.3 mm. Small sample disks (diameter 5 mm.) were taken out of the liner layer, the pipe body and the outside resin layer, to determine the carbon black content by TGA analysis. Some relevant results are listed in Table 5.7 and plotted in Figure 5.27.

Table 5.7 Results of the specific volume determinations on WAVIMAR anti-static pipe at 23~ and an electrical field strength of 4000 V/m.

, , ,, .... ,

WAVIN/~ a n t i - vo lume static pipe, sample set

i

i, total pipe 2, t o t a l p i p e

~,,

2, no resin outside layer

2, liner layer and pipe body

reference system, no carbon black

. . . . . ',, I~,,,~ ,, , ~,

resistivity radial direction,

Ohm.m ,

7.5E5 8.4E6

9.3E6

1.4E7

volume resistivity axial direction,

Ohm.m

2.4E3 2.0E3

2.2E3

3.5E3

3.0El4 , ..... Ull I

a. outside layer- 1.6 %wt. carbon black on resin, b. pipe body �9 I.i %wt. carbon black on resin, c. liner layer - 0.7 %wt. carbon black on resin.

These data show a volume resistivity decrease of about seven decades in the radial direction and a decrease of about eleven decades in the axial direction. The time dependency of the charging currents has completely disappeared but the volume resistivity proves to be field strength dependent. The slight decrease of the volume resistivity from the inner side to the outer side agrees with the increasing carbon black concentrations measured in this direction and this might be an effect of the glass fibre wrapping procedure during the production process.

The WAVIMAR anti-static GFR pipe easily satisfies the electrostatic safety criteria given in 5.3.1, the volume resistivity in axial direction is even too low to protect persons against shortcircuiting. The clear direction dependency of the volume resistivity might be an interesting option for other electric/electronic applications.

2 g

18

18

14

12

16

E E t - O

o 0

X

._> ~

0

E m 0 >

1 8 0

SPECIF IC VOLUME R E S I S T I V I T Y OF WAVIMAR GLASS FIBRE REINFORCED CARBON BLACK F ILLEO EPOXY PIPE IN THE RADIAL OIRECTION AS A FUNCTION OF THE ELECTRICAL F I E L D STRENGTH AND THE P IPE WALL THICKNESS

& ~ WALL THICKNESS 5. I MN (I=IPE AS SUCH) x WALL THICKNESS 4 . 8 NM ( 8 . 3 HM-OUTSIOE-RENQVED) A WALL THICKNESS 2 . 5 MN (2 . 8 MH-OUTSIDE-REMOVED)

&

X

§ x ~

. t ! t i , ! t m B m Ia m ~9 19 m (0 O~ N I.'1 0

Figure 5.27 5 . 8

~/(Field str.ength, E) (Vim) vz ,. ! . I , " i

m IB =a s N N t~ 01

4 . 8

4 . 2

3 . 8

3 . 4

3. g

2 . 6

2 . 2

1 , 8

t . 4

1 . 0 . [ ED Orl

E E

..C 0

X

._> u)

~)

E ,m,. 0 >

SPECIF IC VOLUME R E S I S T I V I T Y OF WAVIMAR GLASS FIBRE REINFORCED. CARBON BLACK F ILLED EPOXY PIPE IN THE AX IAL DIRECTION AS A FUNCTION OF THE ELECTRICAL F I E L D STRENGTH AND THE PIPE WALL THICKNESS

WALL THICKNESS 5 . 2 MM (P IPE AS SUCH) x WALL THICKNESS 4. g NM (B. 3 MH-OUTSIOE-REMOVED) �9 WALL THICKNESS 2. 5 MN ( 2 7 MH-OUTSIOE-REHOVED)

~ X '-'-.-,.-----.-_ ~ X..-.- , , . . ._..__ x

~/(Field strength, E) (V/m) tr~

(It m ,", ~ Ft "q' I/1

181

5.4 Thermally Stimulated Discharge analysis

5.4.1 The TSD techniuue In the.preceding chapters we showed that several thermo- analytical techniques are available to study thermal transitions in polymers. The thermal transitions of new, polymeric systems are usually first investigated by DSC measurements (small amount of sample needed/high scanning speed). The DSC technique is, however, often not sensitive enough to detect weak and/or secondary relaxation effects. DMA (rigid polymers) or dielectric (rigid/rubbery/viscous systems) experiments are then necessary. The sensitivity of the dielectric measurements depends on the polymer's polarisability (see 5.1.2). Besides, DC conduction effects can seriously hamper the detection of the relaxation effects studied. The TSD analysis technique offers in such a case an attractive and sensitive alternative.

TSD experiments are performed using the measuring circuit shown in Figure 5.1. The experimental method is schematically drawn in Figure 5.28. The sample is heated to polarisation temperature (Tp), at time to an electrical field (Eo) starts to polarise the sample. After a certain polarisation time (tpl) the temperature is decreased to To; the polarisation is then 'frozen in'. At temperature To (i.e. at time t,) the electrical field (E~ is removed and a small depolarisation current effect is measured. If this depolarisation effect has been vanished completely, the temperature is to increased linearly as a function of time and the thermally stimulated discharge current is recorded.

5.4.2 Bucci's TSD theory Next to electronic, atomic and orientation polarisation (see 5.1.2), two other polarisation types can occur during TSD experiments: - intrinsic space charge polarisation; free electrons and/or

ions present in the sample, move in the direction of the electrodes and,

- extrinsic space charge polarisation; electrons and/or ions are injected from the outside into the polymer sample.

The electronic and atomic polarisation can not be frozen-in and cause the small depolarisation current after removing of the electrical field at time t2. Thus, the measured discharge current is the sum of dipole and space charge relaxation effects. The exponential relation between the relaxation times of these effects and the temperature makes it possible to shorten the discharge time drastically by therm~l stimulation.

Bucci [36] and Perlman [37] treated these thermally stimulated currents as dipolar relaxation processes with single relaxation times (the Debye model). According to this model, the thermally stimulated current is given by:

T I(T) - A.exp[-E/k.T - B.~exp(-E/k.T) .dT]

To 5.39

"11

r-

gl

c/;

0 Q.

c

N

A Q

.

D 0 V

a -0

0 f ~ ~

0 D

,|m

,.

0

3

m

o I !

-o

-0

r-+

~ ~

~J

183

where: I (T) = the measured TSD current, E = an activation energy, T = absolute temperature, k = Boltzmann's constant.

A = (N.p 2.EO.G) / (k.Tp. T) 5 . 4 0

and

B - i/(E. T) 5.41

where- Eo = electrical field, p = dipole moment, N = number of dipoles, T~ = polarisation temperature,

= characteristic relaxation time T(T) = T.exp(E/k.T), and

E = heating rate.

The shape of the TSD curve is mainly determined by heating rate K, the characteristic relaxation time T and activation energy E. The total amount of charge released is linearly related with field strength Eo. The discharge current is going through a maximum at temperature Tm given by:

Tm = [ (E/k) .K.T.exp(ElkTm) ]0, 5.42

Hence, Tm is not influenced by polarisation temperature Tp nor by the strength of electrical field Eo. The low temperature tail of the TSD curve is described by the first exponent of 5.39 i.e. :

Ln[I(T)] =_ Ln[A] - E/(k.T) 5.43

Thus E follows from the low temperature slope of the Ln[I (T)] versus the inverse, absolute temperature curve.

The theory describing space charge depolarisation processes is complicated [38]. Equations like 5.39 describe the thermally stimulated discharge currents but contain more unknown variables which are depending of- - the type of charge carriers (ions or electrons), - the different possibilities to move (drift or diffusion), - the recombination and dissociation processes during the

polarisation process. TSD dipole relaxation processes are in general easier detectable than space charge depolarisation effects. For non- polar polymers, however, the TSD effect is depending on the space charge polarisation possibilities.

184

5.4.3 Results of TSD experiments _

An example of a PVC orientation depolarisation effect, measured with a combined TMA/TSD system is given in chapter 6. These orientation depolarisation effects were measured on small (i.e. 8 mm.) diameter, samples. Such samples proved to be too small, however, to detect the space charge depolarisation effects in non-polar SSBR rubbers. These non- vulcanised rubber samples were pressed, therefore, at 140~ between two (i mm thick) brass disks with a diameter of respectively Ii0 mm (high potential electrode) and 80 mm (low potential electrode) to a sample thickness of about 0.2 mm. A ring (inner/outer diameter 75/85 mm) of 50 micron thick Vespel foil avoided shortcircuiting between the two brass disks.

The investigated SSBR samples were medium vinyl SSBR systems (BR part about 50 %wt. vinyl BR; styrene content 23 %wt.). The styrene monomer was mainly added during the last stage of the polymerisation process. This resulted in tapered SSBR systems with endblocks i.e. 'tails' with a high styrene content. Detection of these high styrene content tail-structures by DSC or DMA failed. Vanderschueren [39] showed that the space charge depolarisation maxima of SBR systems are structure dependent i.e. styrene content dependent. Hence, the TSD technique was used to investigate these SSBR systems with and without 'tail' structures.

The TSD (space charge) depolarisation currents of three of these non-polar polymers are shown in Figure 5.29, the measured maximum temperatures are listed below-

WRC10805, medium vinyl SSBR - very broad 'tail', Tm = I02oC WRCI0804, medium vinyl SSBR - broad 'tail' , Tm = 82~ WRCI0801, medium vinyl SSBR - standard 'tail' , Tm = 76~

These results confirmed Vanderschueren's conclusion that structural differences are detectable by space charge depolarisation current analysis.

0 PO

0

3 o

r r 0

Q_

r C~

0

C; 0 0 bO

0

0 b 0

0 k~ ,o

TS

D

curr

en

t,

x 1

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mp

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0

0 0

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. :~

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0 0

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0

+ \+

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_.,.

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~ II

,.+

II I'0

II

II g

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3 ..

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"+'--

.~+~

+ j+.

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186

References

i. C.C. Ku and R. Liepins- Electrical properties of polymers, chemical principles, Hanser Publishers, Munich (1987) .

2. W.M. Groenewoud, J. Oil Col. Chem. Assoc., /22, (1979), p. l0 - 17.

3. IEC 93, Recommended methods of test for volume and surface resistivities of electrical insulating materials, the International Electrotechnical Commission, Geneve (1958) .

4. B.V. Hamon, Proc. Inst. Electr. Engr., 99, part IV, (1952) .

5. A. yon Hippel, Dielectrics and Waves, Wiley, London (1954) .

6. N. Hill et.al., Dielectric Properties and Molecular Behaviour, Van Nostrant, New York (1969).

7. N.G. McCrum, B.E. Read and G. Williams, Anelastic and Dielectric Effects in Polymeric Solids, John Wiley, London (1967) .

8. IEC 250, Recommended methods for the determination of the permittivity and dielectric dissipation factor of electrical insulating materials at power, audio and radio frequencies including metre wavelengths, the International Electrotechnical Commision, Geneve (1969) .

9. A. Osier, Z. Angew. Phys., 20, Heft 5, (1966), p. 375. i0. A. Oster, Z. Angew. Phys., 23, Heft 2, (1967), p. 120. ii. G. Weber, Colloid & Polymer Sci., 256, (1078), p. 923. 12. P.J. Phillips, J. Pol. Sci.: Pol. Phys. Ed., 17, (1979),

p. 409. 13. J.D. Ferry, Viscoelastic Properties in Polymers, J. Wiley,

New York (1980) . 14. R.D. McCammon and R.N. Work, Rev. Sci. Inst., 36, (1965),

p. 1169. 15. L.C. Corrado and R.N. Work, Rev. Sci. Inst., 41, (1970),

p. 598. 16. J.C. Coburn and R.H. Boyd, Macromol., i~, (1986), p. 2238. 17. E.B. Murphy and W.A. O'Neil, SPE Journal, 2, (1962), p.

191. 18. R.J. Bender, Handbook of Foamed Plastics, (1965), p. 160. 19. P.A.M. Steeman et.al., Polymer, 32, (1991), p. 523. 20. K. Kadotani, Composites, october (1980), p. 199. 21. G. Banhegyi et.al., Coll. Pol. Sci., 266, (1988), p. 701. 22. A.R. Bunsell, Reinforced Plast., 3, (1984), p. I. 23. J.D. Reid and W.H. Lawrence, J. Appl. Polym. Sci., 31,

(1986), p. 1771. 24. M. Continaud et.al., J. Mat. Sci., i/, (1982), p. 867. 25. Illers (water effect in nylons) 26. D. Brasher, J. Appl. Chem., ~, (1954), p. 62. 27. M. Woo and R. Piggot, J. Compos., 10, (1988), p. 16. 28. Shell Safety Committee, Static Electricity, Shell

Internationale Petroleum Mij. B.V., The Hague, 1976. 29. Silec, Catalogue Generale, (1976). 30. H. Haase, Statische Elektrizitat als Gefahr, Verlag Chemi

GmbH, Weinheim (1972) . 31. Richl. Nr. 4 der Berufsgenossenschaft der Chemische

Industrie, Statische Elektrizitat, Verlag Chemie GmbH, Weinheim, Neufassung (1971) .

187

32. L~ Burton et.al., Rubber Chemistry and Technology, Vol. 62, (1989) , p. 838.

33. J.C.M. Brokken-Zijp, A. Noordam, W.M. Groenewoud and C.H. Klaren, European Patent Application, EP 0 370 586 A2, (1989).

34. H. Kawamoto, Carbon Black-Polymer Composites, The physics of electrically conducting composites, Dekker, New York, (1982), p.135.

35. M.A.J. Michels et.al., Physica A, !57, (1989), p. 529. 36. C.Bucci et.al., Phys. Rev., 14_~, (1966), p. 816. 37. M.M. Perlman, J. Electrochem. Soc.- Solid State Sci. and

Techn., July (1972), p. 892. 38. J. Vanderschueren et.al., Thermally Stimulated Relaxation

in Solids, H4, Field Induced Thermally Stimulated Currents, Topics in Applied Physics, Vol 37, Springer Verlag, New York, (1979) .

39. J. Vanderschueren, Macromolecules, 13, (1980), p. 973.

COUPLED THERMAL ANALYSIS TECHNIQUES

CHAPTER 6

188

CHAPTER 6: COUPLED THERMAL ANALYSIS TECHNIQUES

6.1 Introduction Coupled or simultaneous thermal analysis techniques are gaining more and more importance for several reasons- - using the same sample for two or more analytical

determinations prevents sample difference effects and ensures an identical thermal sample treatment,

- confimation of, or complimentary information obtainable on results, using a second independent technique may justify economically the cost of application of coupled techniques,

- the often small amounts of sample available during process research on new or modified polymeric systems force the use of combined techniques to obtain as much as possible information from a minimum amount of sample,

- new techniqual developments promote the commercial availability of coupled thermal analysis techniques.

A number of combinations with the thermobalance (TG) as basic unit has become wellknown i.e. the TG/DSC or TG/DTA, the TG/MS and the TG/FTIR, the TG/GC/MS, but also DMA/DETA and TMA/DETA combinations have been reported [i - 5]. Coupled thermal analysis techniques usually recall pictures of sophisticated, complicated experimental systems, but also simple combinations are possible. McCammon, Corrado and Coburn [6 - 8] reported many years ago already their dielectric measurements (DETA) in combination with linear thermal expansion (TMA) measurements. A capacitive sample thickness measurement in a dual measuring cell (as described in 5.1.7) provided the information used to calculate the thermal expansion as a function of the temperature in combination with the dielectric data. A still more simple combination of thermal expansion (TMA) and thermally stimulated discharge (TSD) is described in chapter 6.2. A significantly more sophisticated combination of a thermobalance with FTIR and MS coupled in parallel is described in chapter 6.3.

189

6.2 Simultaneous TSD/TMA measurements

6.2.1 The TSD/TMA system The TSD/TMA combination is a typical example of a system developed to measure the same physical property i.e. the glass-rubber transition, on the same sample and at the same time using two independent techniques.

An 'old' but still properly working Perkin Elmer TMS-I was adapted to perform TSD/TMA experiments. The TMS-I is schematically drawn in Figure 6.1 together with some of the important dimensions. The sample, between the probe and the quartz glass sample holder is placed in the furnace. The furnace temperature is programmed to increase linearly with the time. The thermal expansion of the sample is measured via the probe by the linear variable displacement transducer (LVDT). A thermocouple, placed as close as possible to the sample is giving the sample temperature information.

The sample holder system had to be modified to perform TMA/TSD measurements simultaneously. This modification is reproduced enlarged in Figure 6.2. The TSD/TMA sample disk (i mm. thick, 8 mm. diameter), provided with vacuum evaporated silver electrodes on both sides, is placed between a cup-shaped silver high potential electrode and a disk-shaped silver low potential electrode. The TMAprobe rests in the cup-shaped high potential electrode. Both silver electrodes are connected via 0.I mm. thick, glass fibre insulated wires with two BNC connectors in a tufnol holder clamped around the upper part of the sample tube.

The main problem proved to be the mechanical stability and the electrical screening of both electrode wires. Small, thin ceramic pipes proved to give the wires sufficient mechanical stability and electrical insulation. A thin (0.i mm.) brass pipe around both ceramic pipes and the quartz glass sample holder provides the necessary electrical screening. Both electrodes were connected with a voltage supply and an electrometer as shown in Figure 5.1.

Polarisation voltages up to 3200 Volt were applied without any problem. The resulting discharge currents, as low as 1.0xE-13 Ampere, were measured using a Keithley 616 (autoranging) electrometer. The thermal expansion effect measured is the sum of the sample effect and that of both silver electrodes. The thermal expansion of both electrodes (in total 2 mm. silver) was measured between -130oc and 130~ A constant linear thermal expansion coefficient of I.SE-5/K was measured (the Handbook of Chemistry and Physics gives an average value of 1.9E-5/K). This value for the thermal expansion coefficient of silver is, depending on the circumstances, a factor 2 to 20 lower than that of polymer samples investigated. The overall expansion effect measured is corrected for this silver effect if the linear thermal expansion coefficient of the polymer sample as a function of the temperature is required.

190

Sample holder: made from quartz glass, inner diameter: 9.5 mm

Probe: made from quartz glass, probe diameter: 4.0 mm

Oven: inner diameter: 17 mm

Sample length: maximal: 10 mm

Temperature region: -150~ to 325~

Figure 6.1 The Perkin Elmer TMS-1

191

i

i i

probe , -

sample

low potential electrode

I I I I I I I I I

I I I I

, ~ "

I11 A p ~ 1 v . . . .

scale 1 "0.2

=-.-- oven ceramic insulatlon

._-.-, connecting wire

high potential electrode

sample holder

F igure 6 . 2 T h e T S D / T M A s a m p l e ho lde r

192

The thermal expansion of the silver electrodes is neglected if only the temperature location of the glass-rubber transition of the sample is measured.

6.2.2 TSD/TMA results Figure 6.3 shows the results of a simultaneous TSD/TMA measurement on a PVC sample. The glass-rubber transition of this sample is clearly shown by an orientation depolarisation current (see 5.4.2) maximum at 75~ The slightly positive currents before and after the thermally stimulated orientation depolarisation effect are an indication that next to dipole orientation also space charge polarisation effects are involved. The dilatometric measured Tg-value (the transition from the glassy state into the rubbery state is accompanied by a strong increase of the thermal expansitivity) is 76~ The close agreement of the Tg-values measured by the TSD and by the TMA technique indicate that the hypothetical measuring frequency for both techniques is about the same i.e. f(h) is in the order of ixE-2 Hz. to ixE-4 Hz. [9].

The TSD experiments on this PVC sample were, subsequently, repeated at three different polarisation voltages. The resulting depolarisation currents are shown in figure 6.4. Base-lines were drawn as good as possible and the total charge released, determined by integration, is plotted as a function of the electrical field strength in the insert of Figure 6.4. The practically linear relation found between the released charge and the electrical field strength is in agreement with Bucci's theory given in 5.4.2.

This TSD/TMA combination proved to be a convenient (small sample size) and sensitive system for the determination of glass-rubber transition effects of several experimental polymer systems, especially in cases where the DSC technique failed due to a lack of sensitivity. The advantage of this dual technique was especially felt during the investigation of experimental samples with unknown Tg-values. In such situations the confirmation of the Tg-value by a second independent technique is often very valuable.

The sample diameter /thickness ratio (8.0 for this system) proved to be not high enough, however, to perform measurements on non-polar polymers where the TSD effects depend on the much weaker space charge polarisation effects. In such cases the TSD and the expansion coefficient measurements were performed using standard equipment with a TSD sample diameter/thickness ratio of about 400, as decribed in chapter 5.4.3.

193

Figure 6.3 Simul taneous T S D / T M A measurement on a PVC sample with 10%wt. impact improver

+ TMA A TSD resu l t s resu l t s

1.0200

1 . 0 1 5 0

E E 1.01 O0 ..6

c: ,,=.,=.

,,=,.==

E r 1 .0050

Tp = 1 2 0 ~ To = - 1 3 0 ~ tp = 10 mirl

Eo = 1 0 0 0 kV/m sarrl31e thickness = 1.0 mm

sample diam. = 8.0 mm heating rate = 4~

1.0000

0 . 9 9 5 0 20

I , I - , , , I

4O 6 0

I - - . I . . . . . I

8 0

/ .+

- 3 . 8 0

- 3 . 3 0

- 2 . 8 0

- 2 . 3 0

- 1 . 8 0

- 1 . 3 0

- 0 . 8 0

- 0 . 3 0

0.20 100

Temperature, ~

=e Q. E

q l ' "

I I I I

X

c

a

194

Figure 6.4 TSD/TMAsystem: I(TSD) versus E o for a PVC sample with lO%wt, impact improver

+ 1 0 0 0 A 1 5 0 0 0 2 0 0 0 k V/m k V/m k V/m

- 1 2

-10 ~ " 12

- 8 -'~

- 6 - :~" ,~ . . . .

Released charge I Field strength relation

LU X r

~ - 4 o

1:3 or) I -

- 2

1000 1 ( i 00 ~ ILNIO0

Field strength, kV/m

0

2[ 45

| . . . . . . . |

55 II N I . . . . I

65 I

75 , I I

85 95

Temperature,~

195

6.3 The TGA- coupled - FTIR/MS technique

~.3.1 Introduction TGA experiments on polymeric systems often show complex TGA mass/temperature curves in which multiple decomposition products correspond with the weight change observed (see, for example, Figure 2.10). TGA has thus proven to be an excellent quantitative technique but less suitable for specification. This drawback can be eliminated if the components which are causing the mass losses detected, are also analysed simultaneously, the so-called evolved gas analysis (EGA). Several TGA-EGA systems are described in literature, analysing the evolved gases with different techniques i.e. thermal conductivity, cold-trapping followed by GC, mass spectrometry (MS) and infrared (FTIR). MS and FTIR have proven to be the most powerful techniques [3, i0].

User-friendly combinations of TGA/MS and that of TGA/FTIR with the essential reliable coupling of the various system components, became commercially available from different manufacturers in 1987/1988. Both systems have their own strong points. The MS technique is very sensitive and determines both polar and non-polar components, but is mainly qualitative. The FTIR technique, only detecting components with changing dipole moments, is better suited for quantitative determinations. Besides, the FTIR equipment normally used to measure gas-phase spectra, can also be used in the diffused reflectance mode. This offers the additional possibility to investigate a thermally treated sample i.e. a TGA residu.

However, the results of one single detection technique are often not decisive enough to identify a certain component. A TGA coupled with both FTIR and MS (in parallel) should be a much stronger combination. Such a combination was, however, in 1990 not yet commercially available from one single manufacturer.

The need for a quantitative working TGA - coupled - FTIR/MS system was clearly felt in our laboratory: mainly to study the first stage(s) of polymer degradation processes but also to determine small amounts of residual solvent and/or residual monomer(s) in polymeric systems. First, a Perkin Elmer TGA - coupled - FTIR system was purchased and adapted to allow quantitative determinations. A Balzers MS was subsequently purchased and coupled, using the method described by Dufour and Raemaekers [4], in parallel with the FTIR resulting in a TGA - coupled - FTIR/MS system [Ii]. The TGA/FTIR and the TGA/MS coupling, the systems' modifications and calibrations and some typical results are described in this chapter.

196

~,3.2 The TGA/FTIR and TGA/MS coupling The original Perkin Elmer 'vertical furnace' TGA is schematically shown in Figure 6.5. The glass ball-joint coupling offers the opportunity to move the furnace tube downwards (hydraulically) in order to put a sample into the platinum sample pan. The helium purge gas flow (60 ml/minute) is simple and shown in Figure 6.5.

The glass ball part of the ball-joint is replaced by Perkin Elmer to mount a stainless steel capillary lined with a PTFE (Teflon) tube, see Figure 6.6. The evolved TGA gases are sampled just above the sample pan and pass down the heated PTFE line to a 5 cm long, single pass and heated FTIR gas- cell.

The use of a PTFE inner-liner (inner diameter 1.15 mm) allows simple and rapid replacement upon fouling. The total volume of the transfer line and the gas-cell is about 5.8 ml (gas-cell- 4.8 ml). This results in a TGA/FTIR transfer time of about 3.5 seconds. The temperature of the transfer line and the gas-cell are continuously variable between 20~ and 230~ The heated gas-cell is fitted with spring loaded KBr windows to maintain a gas seal at all operating temperatures.

The TGA furnace tube was modified by Balzers according to the ideas of Dufour and Raemaekers [4] to connect the MS heated transfer line, see Figure 6.6. The MS interchangeable capillary tip (see Figure 6.6 insert) is shifted as close as possible to the TGA sample pan. This tip forms the end of a fused silica capillary (inner diameter 0.25 mm) which is connected with the MS and can be heated up to 400~ The FTIR capillary tip/TGA sample pan distance remains constant when the furnace tube is moving downwards to open the system for the sample loading procedure. The MS capillary, however, is connected with the moving part of the furnace tube. Hence, the MS capillary must be positioned in such a way that this tip does not hit the TGA sample pan during the downwards movement of the vertical TGA furnace.

The total gas purge rate was increased from 60 ml/minute to 100 ml/minute, a balance purge of 50 ml/minute, a sample purge of 25 ml/minute and a MS capillary purge of 25 ml/minute (see Figure 6.6). The helium for the MS capillary purge is flowing between the fused silica capillary and its stainless steel support line (see Figure 6.6 insert) from the MS into the TGA. This ~eated gas stream has to keep the MS capillary tip warm enough to avoid condensation at this critical spot. About 5 ml/minute of the total purge gas flow is leaving the system via the MS capillary; the remaining 95 ml/minute is leaving the system via the FTIR capillary.

197

balance purge 40 ml/minute

anti-static shield sample purge 20 ml/minute

ball-joint -

hang-down wire

sample pan (platinum)

TGA oven with the sample temperature sensor

purge gas outlet = - 60 ml/minute

glass furnace tube

Figure 6.5 The Perkin Elmer "Vertical Furnace" TGA 7 system . _ ,

198

FTIR heated transfer line r . . . .

MS heated transfer line

_ . . . . , _

II lm~==..~. ,. m

- ~ _ ~ i : ~ ---7

interchangeable i capillary tip

L sample purge 25 ml/min

-- fused silica capillary, MS inlet, 5ml min

balance purge [ 50 ml/min

' \ ~ I f i . . . . t ~ l l I,a 1 ~ , !ll

I .J-~=; II J ,,

i " i

i �9 . . . .

aluminium MS transfer line support ! I' 'J ' I

I . . . . J

Figure 6.6 The TGA 7 with the flexible, heated TGA-MS and TGA-FTIR transfer lines

FTIR modu 1 �9 1760-X

,,

t r ! g g e r ! 720-VDU u n l t -,

FT IR/HS J manual

8 t a r t

FTIR/HS TRS-?

so'Ptware cont, ro 1 led

s tar t ,

F'TIR opt t e#L1 m o d u l e t ? G O - X

' p P t n t o r ~ - FX-80 |

Tandon 38GSX2B c:omputer

B a I z e r 8 QMG-420 HS

N DTGS de tec to r

" - - " - - I

i ! u s - c e 1 1 ] , ,

i DTGS - : de t .

, .... , ,, , ,, _

HP Co l o ~ PE 77e(~ TRC? p lo t t , e r ~ - 1 computer

,

_ CDS-3 TRS-? L =oft, so~t, i W & P O W & P O 1

~ 0

~ 0

Figure 6.7: Schematic diagram of the TGA-coupled-FTIR/MS system

200

The TGA - coupled - FTIR/MS system as such is schematically drawn in Figure 6.7. The Perkin Elmer 1760-X FTIR is a single beam improved Michelson interferometer with a multicoated KBr beamsplitter giving a wavelength range from 370 to 7200 cm(- i). The heated FTIR gas-cell and a second DTGS IR detector are mounted on an auxilary bench next to the 1760-X FTIR. The primary sample compartment has been left available for 'non- TGA' work i.e. in our case diffuse reflectance measurements.

The Balzers QMG-420-180H MS is a quadrupole analyser with a 1 - 512 a.m.u, mass range and a crossbeam ion source with two filaments. The whole system is controlled by three independant computers during an experiment. The TGA is controlled by the PE-7700 computer using the TAS-7 software. The FTIR spectra measured by the 1760-X FTIR are stored in the 1720-VDU spectroscopy terminal. The MS spectra measured are stored in the Tandon 386SX20 computer. The only 'hardware' connection between these three systems is a unit which triggers the 1720- VDU terminal and the Tandon 386SX20 computer to start at the same time with measuring and storing IR respectively MS spectra. This start can be software controlled from the TAS-7 software or manual.

6.3.3 The heated capillaries tip temperatures Condensation effects in the beginning of the teflon FTIR capillary and in the glass MS interchangeable capillary tip during the first experiments indicated that the temperature in both capillary tips was lower than that in the heated part of the capillaries. Thin thermocouples were mounted subsequently, on these tips to measure the actual temperatures during an experiment. The measured FTIR/MS capillary tip temperatures at different TGA furnace temperatures have been collected in Table 6.1.

Table 6.1 The FTIR/MS capillary tip temperatures as a function of the TGA furnace temperature (FTIR/MS heated transfer lines at 200~

. II I 'r,,, I

~- 'I~A f u r n a c e temperature,

R ~

Ill 100

I 150

20O . ,,, ,,

FTIR capillary tip temperature,

oC

61 ,,

74 ,.,

92

MS capillary tip temperature,

oc

40

55

73

106 92

The results listed in Table 6.1 were disappointing. It was clear that the heating capacity of the helium sample purge via the MS heated transfer line proved insufficient to compensate for the heat losses in the non-heated end-part of the MS capillary.

201

An additional external heating source was used to solve this problem. Two Osram Xenophot HLX 64 635 (15 Volt, 150 Watt) IR heaters mounted on moveable support arms were used to heat the MS capillary tip, a third one was used to heat the FTIR capillary tip. These three heaters are controlled by one Eurotherm 808 controller with the measuring and the alarm thermocouples mounted near the small ball-joint of the MS heated transfer line/TGA furnace coupling, see Figure 6.6. The measurement of the FTIR/MS capillary tip temperatures was repeated using only these extra heating sources (no TGA furnace switched on), the results are listed in Table 6.2.

Table 6.2 The FTIR/MS capillary tip temperatures as a function of the external heaters controlling temperature (FTIR/MS heated transfer lines at 200~

. . . . . , . . . . . . . - , , , , , , ,,

I IR heaters controll temperature,

oc

~ 150

. . . . . ~ ,~,, ,, ',, . . . . . ,,

FTIR c a p i l l a r y MS c a p i l l a r y t i p tip temperature, temperature,

~ ~ 1

7a i 113

, ,

163~ I 215

52

105 ,,

159

212

These data show that the capillary tip temperatures after this modification can be brought in line with the TGA furnace temperature (and hence the sample temperature) up to about 210oc/215oC.

6.3.4 Singl~ com_Don~nt calibration The linear relation between the IR absorption and the sample concentration makes calibration of the TGA - coupled - FTIR possible, for the determination of the total amount of components released during a TGA experiment. Quantitative MS component determinations are more complicated (due to more parameters) than quantitative FTIR component determinations but it might be possible if at least a part of the calibration curve is linear.

The limitations for quantitative determinations of single components (both by the FTIR and the MS) were investigated by measuring calibration curves for the following pure component s-

-n-tetradecane (boiling point 252.5oc, apolar), - benzoic acid (boiling point 249.0~ polar), - glycerol (boiling point 290.0~ polar).

100.0

gO.O

QO.O

70. 0

80.0

441,,0

30. O

211,0

n.tetradecane, 5 mg �9 . m . - " ~ ,.=,, q = , = . m . , m = . ~ / /

A

\ , \

"\

\ \

. c

m o

I !

I !

I !

I !

t !

I !

! f

0 , 0

- -2. 0

_= - -4 . 0

Q a . . . . ~ o

--10. 0 t

I !

- - 12 ,0

DO 0

10.0

o.0 11.0

Time (minutes)

Ss 0 lO~ 0 IS. 0 2O. 0 ;5 . 0

Figure 6.8 The evaporation of n-tetradecane as measured by TGA (mass loss) and DTGA (mass loss rate)

30.0

4.0

203

O. OglO

A

O, O884

0.0488

o . o m e

I1. 0188

O. 0000

I " . . . . . I

G a s - p h a s e s p e c t r u m o f n - t e t r a d e c a n e

r

m c m -1

Figure 6.9A The FTIR spectrum of n-tetradecane vapour

lg00

- I

1000 dl,q0

11,1OA I-

~ 0 g -

11.08 -

0.04 -

o . m -

0.oo

. . . . . . w . . . . . . . . . , . . . . i . . . . . . i A b s o r p t i o n a t 2 9 3 3 c m -1

i"" "' "

==

u

e l l

e=

I /

0 S 1o 18 81 a No

T i m e ( m i n u t e s ) Figure 6.9B The intensity of the CH2 vibration absorption of n-tetradecane vapour as a function of the measuring time

204

Five n-tetradecane samples with sample weights between 1 mg. and 8 mg. were heated from 30~ to 200~ (rate 5~ in the TGA using only the three external IR heaters. The TGA mass/time curve in Figure 6.8 shows that the n-tetradecane sample completely evaporated during these experiments. No traces of condensated product were visible in both capillary tips. Figure 6.9A shows a FTIR gas-phase spectrum with a strong CH2 vibration at a wavelength of 2933 cm(-l) measured during these experiments. The absorption intensity of this vibration is plotted as a function of the experiment time (i.e. the temperature) in Figure 6.9B. The shape of this FTIR intensity/time curve agrees with the TGA first derivative curve, see Figure 6.8. Subsequenly, the FTIR intensity/time curves determined in this way were integrated. These integral values plotted as a function of the TGA sample weights resulted in a TGA - coupled - FTIR calibration curve for n- tetradecane.

Figure 6.10A shows a MS spectrum of n-tetradecane measured during these experiments. It is a typical 'linear alkane' spectrum with the highest m/z-value of 198 from the molecular ion of n-tetradecane and a fragments spectrum with m/z-values differing one CH2 group and an abundance maximum around the C3 and C4 fragments [12]. The intensity of fragment m/z = 57 is plotted as a function of the (experiment) time in Figure 6.10B. The shape of this MS intensity/time curve agrees also with the TGA first derivative curve, see Figure 6.8. Subsequently, the MS intensity/time curves determined in this way were integrated. These integral values plotted as a function of the TGA sample weights resulted in a TGA - coupled - MS calibration curve for n-tetradecane.

Calibration curves for (polar) benzoic acid were measured in the same way without any problem. Glycerol, however, was measured without any problem with the TGA - coupled - FTIR system but the strong polar vapor was not able to pass the tip region of the MS inlet capillary. The thus measured three FTIR calibration curves are shown in Figure 6.11A. These curves for the 'high' boiling point components proved to be linear over the whole concentration range investigated with only small deviations from zero for a sample zero concentration. A series of TGA - coupled - FTIR calibration curves measured on 'low' boiling point components, see Figure 6.12, confirmed the systems' linear behaviour.These results agree with the experimental results reported by Mittleman [13] for C02, S02 and NH3. Figure 6.11B shows the two MS calibration curves measured for n-tetradecane and benzoic acid. Both curves show a linear region for the lower concentrations only. These lower concentrations, however, cover the most interesting region for our area of investigation. The MS calibration curve for water, see Figure 6.13, shows that this linearity also holds for 'low' boiling point components. Hence, a careful TGA sample size choice is making quantitative determinations possible both with the FTIR and the MS for non-polar and polar components with boiling points up to about 250~

eLU!I 6u!JnseeLu eql jo uo!~ounj e se Lg = Z/IN ~ueuJ6mj SIN JnodeA eueoepm~e~,-u jo/q!sue),u! eql

EIO~'9 eJn6!_-I

[u!w] 0C 0~ 0! 0 il I I I I ,inn, , , , i | ' ' ' ....... i J

,,,,-4

Z -4 r

~

fTI !

jnodeA eueoepm]e~-u jo LUm~oeds SIN eql V0 ~'9 eJn6L-I

0000'0

000;'0

000Z '0

000r

000~'0

000g'0

0009'0

000L '0

0008 '0

0006'0

0000' ;

00~ 08T 09T O~T O~T 00T 08 09 0~ 0~ OT-]

,60-3

�9 L0-3

~0~

1.00

T G A - c o u p l e d - F T I R

c a l i b r a t i o n c u r v e s + benzoic & n - t e t r a 0 glycerol

acid decane

0.80

3 r . m

E 0.60

E 4-'

•

E .2 0 .40 Q. 0

.13 <

0 .20

0 .00 [~ 0 2 4 6 8 10

T G A - c o u p l e d - M S

c a l i b r a t i o n c u r v e s + benzoic ~ n - t e t r a

acid d e c ~ e

0 ,45

0 .40

0 .35

m 0.30 <

.E_ 0 .25

X

0 .20 r @ L

0 0 . 1 5

0.10 -

0 .05

0 ,00 0

z~

z~

, I i I , I , , I

2 4 6 8 10

Figure 6 . 1 1 A Sample weight, rag. Figure 6 .11B Sample weight, rag.

.5

.4

.3

.2

. t

0 . 0

o r } (1) ::3

~

E v

E ~

- X r O ~

o J~ <C

~! 02 / N H 3 RL Z BRRT I ON / CRL I BRRT 1 ON URVE / CURVE

x S02 CRt.. I BRRT I ON

A § • ~ M[THRNOL

CRL Z BRRT Z ON ~ + CURVE

X A X

TOLUENE / CRL Z BRRT Z ON

Sample weight (mg)

t"u ~" co oo I$] �9 e - e

Figure 6,12 TGA - coupled - FTIR calibration curves for "low boiling point" components

m ,.4

,.-

f~

I~I

e.e~

;J

eo

~

~ ~,

'

w I

" I

.... w

..

..

.

I

Cur

rent

x

tim

e (A

mpe

re.s

, x

10 -4

)

-.,

~"

8

3-el-

\

4.1t

5. S.~

k

80~

209

6.3.5 Investigation of the thermal decomposition of Cobalt- phthalocyanine by TGA - coupled - FTIR/MS Cobalt-phthalocyanine (CPC), used as a anti-static additive for polymeric systems, is synthesised in two reaction steps- - the reaction of Cobalt(II)phthalocyanine (Figure 6.14A)

with sodium cyanide in a water/methanol mixture to form the precursor, a sodium salt of Cobalt(III)phthalocyanine (Figure 6.14B),

- subsequently 'polymerisation' of this precursor to form a laminar structure where the water ligands (CPC system A, Figure 6.14C) or the ammonia ligands (CPC system B, Figure 6.14D) form hydrogen bonds with the cyanide ligands of the adjacent Cobalt-phthalocyanine rings.

This second step of the synthesis is carried out in water (CPC, system A) or in a aqueous solution of ammonia (CPC, system B) at I00~ Next to ammonia, some water will always 'built-in' as ligand in CPC, system B. With a number of TGA - coupled - FTIR/MS experiments, the amount of ammonia and water ligand in two CPC, system B and in two CPC, system A samples was determined. Besides, the first steps of the thermal decomposition of the CPC, systems A/B were investigated.

The samples were heated, during these experiments, in a helium atmosphere from 40oc to 450~ at a heating rate of 10~ minute. The released vapours were analysed simultaneously by FTIR and MS. The TGA mass/time (M/t) and the first derivative (dM/dt) curves of such an experiment with a CPC, system B sample are shown in Figure 6.15. The sample is losing about six percent of its mass in at least two steps. The FTIR gas- phase spectra measured during such a 'first step' clearly shows specific absorptions of ammonia at 933 and 966 cm-l. The intensity of the most dominant NH3 absorption at 966 cm-I and the intensity of the m/z = 17 (the molecular weight of ammonia) are also plotted as a function of the measuring time i.e. temperature in Figure 6.15. Obviously, This CPC, system B sample looses its ammonia mainly during the first step of the decomposition process.

The FTIR gas-phase spectrum of the decomposition products of a CPC, system B sample measured at 260~ (i.e. after 22 minutes) is shown in Figure 6.16 after subtraction of the dominant ammonia absorptions. The mass spectrum measured at the same time/ temperature is also shown in this figure. This FTIR spectrum shows a few low intensity absorptions at 716 cm-I and at 3272/3337 cm-i next to the CO2 (2337/2365 cm-l) and H20 (4000 - 3500 and 2000 - 1200 cm-l). The mass spectrum shows, besides the m/z values corresponding with ammonia (m/z = 15/16/17), strong m/z = 26/27 components and a smaller m/z = 52 component. The observed FTIR absorptions are hydrogen stretch and bending vibrations of a triple bonded carbon - hydrogen bond. In combination with the observed m/z = 27 value, it was identified as hydrogencyanide (HCN).

210

FIGURE A: STRUCTURE OF COBALT(II)- PHTHALOCYANINE

�9 , . , .

x _1 / ' / Co

I I--

CN

Nil +

X - OH 2 HCIM CH3CHL, OH

FIGURE B: STRUCTURE OF THE PRECURSOR

r //-~) -/7 I - - 0

IT "H''--,NC

eN..---H~N

H,O,s

CN

I

/

FIGURE C: POSSIBLE STRUCTURE OF SYSTEM A

FIGURE D: POSSIBLE STRUCTURE OF SYSTEM B

Figure 6.14 Possible structure and synthesis of cobalt-phthalocyanine

100. 0

~.wt �9

96. 5

93. 0

211

/ d M / d t c u r v e

A / :T~: ~_

, X~ ~~176 ..... .... ," ,

0 5 I0 15 20 25 ~I0 ~5 41

O. 0150 '"

A

O. 0075 -

O. 0000 0

i t Jr ' '~ . . . . . . ! , I -

5 10 ! 5 ZO 25 30 315

9 6 6 c m - 1

41

1,~] 1 2 O O ,

. 1 O O O

. ~ O O

. ~1~OO

. O 2 O O

I N T , R A ~ : [ 1,1='-O6 3

m / z - 17

O 10 2 0 3 0 <10

Figure 6.15 Release of ammonia during thermal degradation of CPC, system B sample

212

0.010

A 0.008

O. 008

O. 004

O. 002

O. 000 4000

I I I I I I = - -

= .

==. . .=

t

~1500 3000 Z500 2000 1500 I000 450

cm"

- 06 -

- 07 .

- (Z~"

- 09 - i , . _ . , _ . . I ~"~2'~~~"~~~~~~~"~"~~'~~"~~~~~"~"~~~~~~~~~~~~+~~~~~~ Figure 6.16 FTIR and mass spectra of the vapours released by a CPC, system B sample at 260~

I00.0

~wt.

g 8 . 5

g3. 0

213

d M / d t c u r v e

M / t c u r v e

~ w l ; . / m f r ' l .

0 5 10 15 20 25 30 35 41

O. OOlO '

A

o. 00o5 -

O. 0000 ~ "

0

I I I I

,| ,,

5 10 t 5 20 25 310

3 2 7 2 o r e - 1

I

35

===t 4 5 ~

. 3 5 0 0 -

. 3 ~ -

�9 2 5 0 0

. 2 ~

. 1 5 O O

. I O O O

. 0 5 0 0

. 0 0 0 0

I N T , R A N G I E : [ 1 E - O 7 ]

m / z = 2 7

O 1 O 2 0 3 0

Figure 6.17 Release of hydrogencyanide during thermal degradation of a CPC, system B sample

4 0

214

The intensity of the absorption at 3272 cm-i and the intensity of m/z = 27, plotted as a function of the measuring time are shown in Figure 6.17. Obviously, HCN is released during both steps of the decomposition process! FTIR and MS spectra measured during the second decomposition step at 290~ (i.e. after 25 minutes) are shown in Figure 6.18. The FTIR spectrum contains the same absorptions as shown in Figure 6.16 (H20, C02 and HCN). In this spectrum, however, two weak absorptions at 2144/2166 cm-i are visible. The mass spectrum shows m/z values of CO2 (m/z = 44), HCN (m/z = 26/27)and a strong m/z = 52 signal. The absorptions at 2144 and 2166 cm-I are cyanide stretch vibrations. In combination with m/z = 52 it was identified as cyanogen (ethanedinitrile, NC-CN).

The intensity of the absorption at 2166 cm-i and the intensity of m/z = 52, plotted as a function of the measuring time are shown in Figure 6.19. The results in this figure clearly show that cyanogen is released mainly during the second step of the decomposition processZ There appears to be a large difference in intensity of the FTIR absorption and the corresponding molecular ion, m/z = 52. Cyanogen is, however, a highly symmetrical molecule, resulting in IR absorptions of very low intensity.

Detection of water and C02 with FTIR at low concentrations is unreliable, due to their presence in the atmosphere. These compounds can be detected with MS. The release of water (m/z = 18) and C02 (m/z = 44) during the decomposition of a CPC, system B sample is shown in Figure 6.20. Absorbed water is released during the first ten minutes of the experiment (40~ - 140~ followed by an amount of chemically bound water at the very beginning of the first decomposition step (the thermal stability of the ammonia ligand is thus higher than that of the water ligand!).

C02 is mainly released during the second decomposition step, and at temperatures higher than 350~ (phthalocyanine ring degradation). The m/z = 18 curve in figure 6.20 shows that the release of C02 (and cyanogen) during the second decomposition step is accompanied by a decrease of the released water intensity. This might be caused by cyanogen reacting at these elevated temperatures with traces of water to form HCN and C02.

Figure 6.21 summarises the results described above. The deomposition products shown in this figure for a CPC, system B sample, were also measured for the CPC, system A samples, only differences in the measured intensities of released water (quantitatively determined with the aid of a MS calibration curve) and ammonia (quantitatively determined with the aid of a FTIR calibration curve) were found-

215

O. 0 0 5 5

o . 0 0 4 4

O. (1033

O. 0 0 2 2

O. 0011

O. 0 0 0 0

4 0 0 0

. . . . i i' ~ ~ i " 't I i

, , I , I I _

3 5 0 0 3 0 0 0 Z 5 0 0 2 0 0 0 1 5 0 0 l O 0 0 4 5 0

c m -1

-06-

-OT-

-08-

09- ."T"'."T""'"I"'"'"T'"m"I""m'T'"I"'T'"I"'T'"""T'"m"r'"""T'"m"r '""1' I, i, '-r"l 20 40 60 80 100 120 140

Figure 6.18 FTIR and mass spectra of the vapours released by a CPC, system B sample at 290~

100. 0

~ w t .

96 .5

\ !\ I r a i n .

93. 0

216

0 5 10 15 20 25 30 35 41

O. 0007

O. D003

O. 13000

- I I ' ' I ' i

i I I I

5 tO t5 20

/

I I I

2 1 6 6 c m - 1

, , li , |

25 3O 35 41

. 1 6 ~ -

. 1 ' q 0 0 -

. 1 2 ~ -

. 1 0 0 0

. ~ : ~

. ~ 2 ~

INT. RAI~K:;E:: [ 1E-O6 ]

m/z - 52

Figure 6.19 Release of cyanogen during thermal degradation of a CPC, system B sample

i .... ,

217

100.0 ~.wt.

g6.5

93.0

dM/dt c u r v e

M/t c u r v e

~W%. / mtn.

0 5 10 15 20 25 ~10 35 41

. 1 6 0 0

. 1 4 0 0

, 1 2 0 0

.1OO0

, 0 8 0 0

. 0 6 0 0

, 0 4 0 0

INT. RANQE: [ 1E-O6 ]

m / z = 18

_ - : ~ _ ~ - ~ : .

�9 �9 ~ . . . . �9 . . . . . . = ~

0 1 0 2 0 3 0 4 0

. 2 0 0 0 l g 0 0

1 6 0 0

1 4 0 0

1 2 0 0

1 0 0 0

. 0 e 0 0

. 0 6 0 0

. 0 4 0 0

. 0 2 0 0

. 0 0 0 0

I N T . R A N G E : [ 1 E - 0 ? ]

m / z = 44

0 1 0 2 0 3 0 4 0

Figure 6.20 Release of water and C02 during thermal degradation of a CPC, system B sample

218

10{3

~ w t �9

9g

98

g7

g6

gS

g4

g3

g l -

go 4O

d H / d % o u r v e

H / t , c u r v e

H20

NH 3

HCN

NC-CN

I I I

I00 150 200

C02 1 I ..... t t

250 300 35O 400

I t o t a l c o m p o n e n t , r e l e a s e r e g t o n

~W% o / m t n .

450

Figure 6.21 Release of the identified products during thermal degradation of a CPC, system B sample

219

Cobalt phthalo cyanide systems

absorbed bound water, bound ammonia, water, %wt. %wt. %wt.

B1 0.4 0.3 1.9 B2 0.3 0.2 2.6

A1 0.7 1.9 0.3 A2 0.6 I.i 1.1

The total amounts of bound water and ammonia are respectively 2.2, 2.8, 2.2 and 2.2 %wt. The theoretical amount, based on the possible structures given in the Figures 6.14C and D, are 2.84 %wt. for only a NH3 ligand and 3.01%wt. for only a water ligand.

Cobalt(II)phthalocyanine (Figure 6.14A) and the precursor (Figure 6.14B) were also measured to exclude the possibility that ammonia is already a decomposition product of these components. No ammonia could be detected during these experiments.

A series of Diffuse Reflection measurements was performed, subsequently, on partly decomposed CPC (system B) samples to check the decomposition pattern given in Figure 6.21. Four samples were (under the same conditions) heated to different, previously selected temperatures, see Figure 6.22. The measured IR Diffuse Reflection spectra of the TGA residues of these four samples offer the possibility to monitor the changes in the molecular structure in an independent way.

The measured spectra are shown in Figure 6.23. Spectrum A is the reference spectrum of the non-treated sample. The presence of ammonia in the molecular structure results in four N-H stretch vibrations between 3100 and 3400 cm-l. The cyanide ligand is visible by two vibrations at 2152 and 2214 cm-l. Spectrum B is measured after heating the sample up to 140~ only causing the release of absorbed water. The nearly identical A and B spectra confirm this conclusion. Spectrum C is measured after heating the sample up to 260~ i.e. in the middle of the first degradation step. The release of ammonia is confirmed by the clear intensity decrease of the N-H vibrations in the TGA residu sample.

There also appears to be in spectrum C a decrease of the cyanide vibration at 2152 cm-i accompanied by an increase of the cyanide vibration at 2214 cm-l. This change in intensity of these two vibrations is completed after heating the sample up to 280~ (spectrum D). The N-H vibrations of the ammonia ligand have now completely disappeared. When the sample is heated, subsequently, up to 350~ (spectrum E), the cyanide vibration has prac~lcally vanlsne~.

220

tO0 . O

~W' t , .

9 6 . 6

g7 . 2

g 5 . 6

g 4 . 4

93. 0 40

I " i .... I ' 'it

d H / d % c u r ' v e

H / % c u P v o

I m ! I ~ , I I ...... !

T e m p e r a t . u r e ( d e 9 . C )

Figure 6.22 Indication of the maximum temperatures after which the DRIFT spectra were measured

~ W t . /

r a i n .

4 5 0

221

1 .50

A

1 .35

l . ZO

1 .05

0 . 9 0

O. 75

0. 80

0. 45

0. 30

O. 15

i I I ' ' i .... I I i I

, 1

0.00 3500 3200 3000 2800 2600 2400 2200

c m -1

Figure 6.23 DRIFT spectra of a CPC, system B sample taken after different thermal treatments

- ~ ~ ~ Spec t rum B ^

S p e c t r u m FI A I

, .,~'I, t . . . . . . . . . . . .

2000 1800 [575

222

The two cyanide vibrations are the result of the presence of 'free' (terminal) cyanide ligands and cyanide ligands that form hydrogen bonds with water/ammonia in these systems. Hydrogen bonding weakens the triple carbon - nitrogen bond, resulting in a shift to a lower frequency. This means that the observed vibration at 2151 cm-i is caused by the hydrogen bonded cyanide ligand, while the vibration at 2214 cm-I is the result of the presence of 'free' cyanide ligand. The shift in intensity observed after the first decomposition step is thus caused by a loss of hydrogen bonding between the cyanide and the water/ammonia ligand. This can also be the reason for the release of some hydrogencyanide during the first decomposition step.

6.3.6 Investigation of the released vapours d~ring .... the cure of an epoxy resin system by TGA - coupled - FTIR/MS. The investigated epoxy resin system consisting of the glycidyl ether of phenol novolack (GEPN) with dicyanediamide as curing agent (added as an diglycidyl ether bisphenol A (DGEBA)/dicy 60/40 masterbatch) and an accelerator in a 100/15/0.5 ratio. + This GEPN/dicy system is used to prepare epoxy resin based castings. The mould, filled with the GEPN/dicy system, has to be heated slowly to the cure temperature of about 170~ to properly release the heat generated by the exothermic cure reaction. Serious thermal degradation effects inside the casting and the release of bad smell and possibly unhealthy degradation products occur if this reaction heat is not released sufficiently. A series of TGA - coupled - FTIR/MS experiments was performed to investigate the conditions resulting in the release of degradation products during the cure of GEPN/dicy systems.

The reaction mechanism of epoxy groups with dicyanediamide is complex [14]. We therefore confine ourselves to say that GEPN and dicyanediamide are forming a solid resin matrix during an exothermic reaction. The reaction heat can cause thermal breakdown of the matrix material just built-up. Earlier performed experiments pointed at the release of ammonia and possibly some pyridine-like components during this thermal degradation process. But the release of water and a phenolic- OH residue can also be expected, see Figure 6.24.

About sixty milligramme of the GEPN/dicy system was heated in the TGA - coupled - FTIR/MS from 30~ to 400~ (rate 5~ to measure straightforward the volatile thermal degradation products of the ~ resin matrix material. The GEPN/dicy sample was cured (according to a DSC experiment under the same thermal conditions) between 160~ and 240~ Detectable mass losses due to thermal degradation processes startea at 295~ {onset temperature). The mass loss rate proved to be maximal at 385~176

- -- O__. CH 2_..._ cH/O_.._~CH 2

~~R~.~ O--CH 2- cH/O"CH 2

H

§ H2N.__C__ N__C~ N

!1 NN

. \ / M R T R I X

M R T R [ X ~ . ~ "- + "~ " <Q OH \

R

NH3 4. H20

Figure 6.24 GEPWdicy resin system, cure and thermal degradation processes (schematically)

E-07 FIGURE B. The mass spectrum at 340 C of the released vapours during the degra-

E-O~ dation experiment of a cured, GEPN /dicy system sample in the TGA

E'e9

E-10' .. ~ . . . . . ,--.-,.--,.,--.-,.,--.-, ....... , .... , . - , - . ,

20 40 60 80 100 120 140

E-07

E-08

E-09

E-10

FIGURE C. The mass spectrum of a tri-methylpyridine reference sample as measured by the TGA- coupled-FTIR/MS system

I - iII-- I ~ -'-T -T -I--'---I I -'- '---I----'-- I-- I -I .... '--I-- "- I --'-- i

20 40 60 80 100 120 1,t0

Absorption &MI A

i.elm n

! = �9 l II | L ~

1179 cm- l t Pheno l i c OIN v i b r a t i o n ~ = / 1 1 7 9 cm-I "

& l l

/ 966 c m - i l N H 3 v i b r a t i o n

�9 Its

L eta

2SS=C =~O'C ~ ~O=C 3SS'C 3aO=C t I - - , I I t - o . m

4s so ss eo es 'Time, min. 7 3

HGURE A. The NH3 and phenolic-OH concentrations versus time/ Temperature as measured by FTIR in the vapours Feleased during the degradation experiment of a cured GEPN/dicy system sample in the TGA.

Figure 6.25 Results of FTIR/MS analysis of the vapours released during the thermal degradation of a GEPN/dicy system sample

225

The FTIR gas-phase spectra are showing clearly the absorption maxima of phenolic-OH components (characteristic vibrations at 1179 and 3652 cm-l), of NH3 (characteristic vibration at 966 cm-l) and that of water. Figure 6.25A is showing the intensity of the characteristic vibrations of the two first mentioned components as a function of the time/temperature. The highest NH3 release rate occured at about 350~ while the release of phenolic-OH components reached their highest level at about at about 390~

Pyridine-like components can not be detected by the FTIR in this case due to overlapping vibrations from the phenolic-OH components. The MS spectra, see Figure 6.25B, are showing next to the NH3 (m/z = 17) and the water (m/z = 18) contributions a clearly present component with a maximum m/z-value of 121/122. Tri-methylpyridine was measured, subsequently, as reference, see Figure 6.25C. The close agreement between both spectra makes it reasonable to assume that pyridine-like components are present in the released vapours during the thermal degradation of a cured GEPN/dicy resin sample.

The degradation onset temperature measured during this experiment (295~ is a dynamic value which is strongly heating rate dependent. A static degradation onset temperature has in fact to be used to indicate the maximum temperature which can be reached in a GEPN/dicy based casting (during the cure process) before thermal degradation effects will occur. Experiments to determine the thermal stability of polypropylene (see 2.2.2) showed that the semi-static degradation onset temperatures measured during ultra-low heating rate TGA experiments were close to the isothermally measured, static degradation onset values. The semi-static degradation onset temperature of the cured GEPN/dicy system measured in this way proves to be 240~ (241~ - 0.2 %wt. mass loss).

Subsequently, the release of ammonia (the only degradation product which could be measured quantitatively) was measured during a series of TGA experiments with variable heating rates. These experiments were performed to show how the heating rate influences the temperature and thus the possible occurance of thermal degradation, in a (small) curing TGA sample.

Non-cured GEPN/dicy samples (about 60 mg.) were heated from 150~ to 250~ with heating rates increasing from 40~ to 140~ Each sample was then kept isothermally at 250"C until a total experiment time of ten minutes was reached. Minor sample mass losses were measured during the experiments with heating rates up to 100~ The higher nearing rate experlments, however, resulte~ in serious mass losses, see Figure 6.26. The amounts of ammonia released during these experiments are plotted as a function of the heating rate in Figure 6.27. These amounts of ammonia were all measured during the first two minutes of each experiment. The presence of other thermal degradation products (water,

9 4 . 0 -

~ g2. o - e -

. . , _

I D N 9o.o

88. 0

100. 0 - ~ _ ~ _

g 8 . 0 188 d e g . C / m i n u t e

15e - 2 5 8 ~ mr. v t r t t b ! e / ' , \ _ h e a ~ t - O r a t e , ~ ~ \

96. 0 + ! s o t h e r ' m t | I t , 25B d e g . O t ( e t c h e x p e r i m e n t t o t a l l y ~ i ~ ~ . _ . ~ l a dog C / m i n u t e

I 8 m i n u t e s ) I \ \

86. 0 -

8 4 . 0 -

82. 0 - 0 . 0

/ ,

i ,

\

1 3 B d e 9 . C / r o t n u t s

[ 4 B d e p . C / m t n u t . e

i i i I T I I I I 0 . 2 0 . 4 0 . 8 0 . 8 1 .0 1 .2 1 .4 1 . 8 1 . 8

T i m e (m inu tes )

Figure 6.26 -he mass losses measured during the cure of GEPN/dicy system TGA samples at different heating rates

- I 0 0 . 0

- 9 8 . 0

- 9 8 . 0

94. 0

- 9 2 . 0

90. 0

- 8 8 . 0

- 8 6 . 0

- 8 4 . 0

- 8 2 . 0 2 . 0

L~ b3 O~

m o

40 ~

5O

6O i

?e ,D

80

--4

98

o :s

II0

ID

.

Im

,.

120

=,

o

130

%

-

3 :3

140

�9

m

O

t~p

fJl

i

Ul 0D

, !

-%1

r ,

,,, .

,, ,

,|,

, O

| i

!

r

| z

~3o

.

dp

c

O (P

CA

I 'L

~0

o o lu

Is

o o.

Z :I:

(r w

228

pyridine-like components and phenolic-OH components) was only detected during the first two minutes of the 130~ and 140~ minute heating rate experiments. The concentrations of these components released during the experiments with heating rates lower than 130~ proved to be too low for a proper detection.

The shape of the curve in Figure 6.27 and the fact that ammonia (and the other degradation products) are produced only during the first two minutes of these experiments are indicating that for heating rates higher than 100~ the real temperature of the sample in the TGA pan was exceeding, during a short period of time, the temperature indicated by the TGA temperature sensor. It is assumed that sample cure and sample degradation processes occured simultaneously for a short period of time during these experiments.

Thus, it proves possible to determine for a small GEPN/dicy TGA sample the rather sharp transition from a non-isothermal cure process without hardly any thermal degradation to a situation where cure and degradation processes occur simultaneously. The critical heating rate of a GEPN/dicy casting system will be strongly casting shape/dimensions dependent. The only way to prevent any thermal degradation during the cure of a GEPN/dicy casting system is to use such a heating rate that the temperature in the casting will not exceed the (static) thermal degradation onset temperature of 240oc.

The examples of the use of coupled TA techniques given in 6.3.5 and 6.3.6 show the extended amount of information which is obtained by the use of such a strong combination of techniques. We are convinced that these experiments offered better results than in the case separated/single TGA, FTIR and MS systems were used. This because we experienced that a proper choice of experimental TGA conditions remained always the first (and often critical) step of a series of experimental steps necessary to reach such a result. In other words, during investigations like this, the TGA technique remains the basic technique, while the FTIR and MS are used as sophisticated and sensitive detection systems.

229

References

i. D. Compton et.al., Research and Development, April, (1989), p. 68.

2. E. Charsly et.al., American Laboratory, January, (1990) . 3. J.P. Redfern, Polymer International, 26, (1991), p. 2819. 4. P.R. Dufour and K.G.H. Raemaekers, Thermochimica Acta,

175, (1991), p. 263. 5. W.H. McClennen et.al., Anal. Chem., 65, (1993), p. 2819. 6. R.D. McCammon and R.N. Work, Rev. Sci. Inst., 36, (1965),

p. 1169. 7. L.C. Corrado and R.N. Work, Rev. Sci. Inst., 41, (1970),

p. 598. 8. J.C. Coburn and R.H. Boyd, Macromol., 19, (1986), p. 2238. 9. P.J. Phillips, J. Pol. Sci.- Pol. Phys. Ed., 17, (1979),

p. 409. i0. J.A.J. Jansen and W. de Haas, Analytica Chimica Acta, 196,

(1987), p. 69. II. W.M. Groenewoud and W. de Jong, Thermochimica Acta, in

press. 12. F.W. McLafferty- Interpretation of Mass Spectra,

University Science Books, Mill Valley, California, (1980) . 13. M. Mittleman, Thermochimica Acta, 166, (1990), p. 301. 14. M.A. Clayton, Epoxy Resins Chemistry and Technology, M.

Dekker Inc., New York, p. 501.

CHEMICAL STRUCTURE / PHYSICAL PROPERTIES CORRELATIONS

CHAPTER 7

230

CHAPTER 7: CHEMICAL STRUCTURE/PHYSICAL PROPERTIES CORRELATIONS

7.1 Introduction The continuously ongoing process of replacing or supplementing the more traditional materials such as wood, natural fibres and metals by synthetic polymers, is stimulating the development of polymeric systems covering either an even wider range or a more specific set of properties. The ability to estimate the chemical/physical properties of a polymeric system from its molecular structure before such a polymer is synthesised can reduce the research costs considerably. Besides, a thorough knowledge of chemical structure/physical properties relations is necessary to improve with success the balance of properties for current commercial polymeric systems. The list of important physical properties of a polymer is, however, long. We mention, without pretending to be complete:

A. phase transitions, volumetric properties, calorimetric properties, electrical/optical/magnetical properties, mechanical and acoustic properties.

B. properties determining mass transfer in polymers, rheological properties of polymer melts, chemical, thermal and UV stability.

C. ultimate mechanical and electrical properties such as" creep, failure, toughness, hardness, friction, wear, yield strength, tracking and dielectric strength. ageing effects.

Many chemical structure/physical property relations have been reported in literature but the most important contributions in this field have been made, we think, by Van Krevelen [1] and Bicerano [2]. Both authors present a total concept for polymer properties/molecular structure correlations based on the group contribution technique (Van Krevelen) and on connectivity indices calculations (Bicerano) covering all the properties mentioned in group A and B and some of the properties of group C. Seitz published a concept to estimate the mechanical properties (from group A and C) of polymers from their molecular structure [3].

It is difficult however to speak about 'the properties' of a polymeric system due to the many variables which determine the ond-ugo vmluoa. Thi~ han~orQ mQrlouol)- thQ uoo of tho estimated or calculated polymer property values. For example, nearly all properties mentioned in group A, B and C are temperature dependent while an important number of these properties is time (and pressure) dependent.

231

Other factors determining end-use properties can be:

- possible tacticity i.e. a-, iso- or syndio-tacticity, - possible cis/trans isomerism, - molecular weight and molecular weight distribution, - copolymerisation/blending, - plasticiser/filler additions, - degree of crosslinking, - effects of different processing techniques, - presence of residual monomers and/or solvents (water).

This extended number of variables prompted us to work in the context of this book only with a set of so-called 'key- properties' instead of trying to calculate immediately all the properties mentioned in group A, B and C. As key-properties are considered:

- the Tg-value, - the Tm-value, - the Hf-value, - the thermal stability and - the moisture sensitivity.

Nearly all properties mentioned under group A are influenced by the temperature locations of the Tg/Tm values and the crystallinity level (Hf-value). The Tm-value and the thermal stability determine together the polymer's processing window. The moisture sensitivity, finally, is important in connection with the barrier properties of a polymeric system. Furthermore moisture absorption influences several physical properties considerably (see chapter 5.2) .

In this chapter an attempt is made to improve some of the existing molecular structure/physical properties correlations. A set of consistently measured TA data is used, in order to estimate the above mentioned ,key-properties, of polymeric systems.

232

7.2 The Tg-value estimation

7.2.1 Introduction The transition region in which an amorphous glassy polymer changes from its glassy state into a rubber-like state is important because dramatic changes in the polymer's physical properties are observed during this transition. These changes are completely reversible as the transition from a glassy state into a rubbery state is a function of molecular motion, and not the polymer structure. The temperature location of this glass-rubber transition region however strongly depends on the molecular structure. This temperature location is usually indicated by one, single temperature: the Tg-value (see 1.1.3).

The Tg-value is, according to Cowie [4], depending on-

- chain flexibility ) - polarity ) molecular structure - steric effects ) effects - tacticity and cis/trans isomery) - branching and crosslinking ) - molar mass, - the presence of additives, fillers, residual monomers

and/or impurities. - morphological effects, especially crystallinity, - rate and type of measurement.

This extended list of factors affecting the Tg-value, is one of the major causes of the widely differing experimental values reported in literature for various common polymers.

In literature, several molecular structure/Tg-value correlations are reported in spite of the often considerable amount of scatter in the reported Tg-values. It is not surprising that many of these Tg-value estimation methods are based on correlations with the chain stiffness and/or the cohesive forces, since chain flexibility is considered to be the most important Tg-value influencing factor. A well known empirical, molecular structure/Tg-value correlation based on the group contribution additivity is given by Van Krevelen [1]. Kreibich and Batzer [5] developed a correlation based on cohesive energy (Ecoh.) values calculated according to the method published by Fedors [6]. Wiff and Altieri calculated Tg-values using the relation between the density and the log Tg [7]. Hopfinger [8] used a molecular modeling technique in combination with a group additive property model to calculate Tg-values of linear polymers. Bicerano developed an impressive, and we do think the most extensive, Tg-value calculation method based on connectivity indices [2]. Without L,~laa~ ~umli~l=L~, ~v~al ULIA~L ~U~ILLIDuLIuII~ UII LIII~ ~UDJ~:~

can be found in [9 - 12] .

233

7.2.2 The 'modified cohesion e~eruy' method We expected that an improvement of the Tg-value estimation was still possible by careful selection of the reference Tg-values used. These Tg-values should preferably be self determined values measured under standard conditions (see 1.1.3) or selected literature values. They should have been determined on amorphous or low crystalline polymers with a sufficient high molecular weight i.e. Mn value a 50.000. Besides, these polymers should not contain fillers, plasticisers, residual monomer and/or solvents.

We tried to improve the Kreibich and Balzers Ecoh./Tg-value correlation. This method was chosen because of a) its simplicity compared with for instance Bicerano's method and b) the possibility to calculate the Ecoh. value of every polymer repeating unit from the ~Ecoh. values of the structural groups given by Fedors in case experimental Tg-values are not available. The latter differs with for example Van Krevelen's method which needs experimental Tg-values to calculate the group contributions. Kreibich and Batzer [5] reported a linear relation between the Tg-value and the cohesion energy (Ecoh.) at 298 K:

Tg-value, (K) = 0.0145 x (Ecoh./Ea, i) + 120 7.1

where: Ecoh. - cohesion energy, J/mol. Ea, i = number of independent moving/rotating groups.

Kreibich and Batzer used the structural group 6Ecoh. data of Fedor [6] to calculate the necessary Ecoh. values. The average (Tgexp. - Tgcalc.) value of 63 calculated Tg-values proved to be _+ 15 K, but for 19 of these 63 values this difference is larger than 15 K (between 16 K and 62 K).

we tried to modify a number of Fedors structural group ~Ecoh. values, using a basis set of 22 self measured Tg-values and some selected literature values, to improve the fit of the Ecoh./Tg correlation. This proved to be possible only if a distinction was made between polymers with only C-atoms in the mainchain and the other, mainly C6-rings and -0- containing mainchains.

Thus, two sets of ~Ecoh.-values for polymer structure elements were calculated, one for linear polymers with five or more succeeding C-atoms in their mainchain and one for the other linear polymers, see Table 7.1. Using the values given in Table 7.1, the Tg-value of linear polymers can be calculated according to equation 7.1 with what we call the 'modified cohesive energy' method.

234

Table 7.1 The ~Ecoh. values of polymer structure elements.

~__ r,,,, ~ , . . . . . , ii ,I _ .' " x ,, , , L H t + , .

structure element ~a, i ~Ecoh., J/mole 8Ecoh., J/mole -(C)n- n ~ 5 + -(C)n-, n ~ 4

. .,,. ~ .... ~ - + , . . . . .

-CH3 0 (CH3-)2, symm. 0 CH3-(C6-ring) 0 CH3-(C6-ring)-CH3 0 -CH3, see note 0 -CH2- 1 !

-CH- 1 I

-C- I

-CH= (cis) -CH= (trans)

i I

-C= (cis) -~= (trans)

-C6H4-/-C6H5 -C6H4-C (CH3) 2-C6H4-

-O- -CO-

-OH - N H -

9826 5040 5040

- 8617 - 17234 - 2928

4936 4936

3430

1464

1129 3196

3430

1464

z/o 3

2878 3982

27123

3346 18405

23747 54839

27951 85007

4852 29243

28708

-CmN -SO2-

26503 - 42064

-C1 in-HCC1- -C1 in =CCI-, cis (-Cl} 2, symm.

: C1 - (C6 - ring) C1- (C6-ring) -C1

-Br Br- (C6- ring) -Br

-F (-F} 2, symm.

- N -

- S -

23885 13176 11922

- 13176 - 16230

0 i 0

1 1

15477 15477 - 21403

18372 6132 9169

4183 4183 - 5566

-C~HI 0-/-C~W11 -(CO)-O-CH3 in PMA and in PMMA -0- (CO) -CH3 (PVAc)

I/• ?~AK4 0 14117 0 20188 0 16875

1 I I , , I ,,,l,,l, '" " , , , , , , , , , ' '" =

-CH3 in [cn-cH~3-0]m-, n ~ 3.

235

This 'modified cohesive energy' method does not permit the calculation of Tg-values of polymers with bulky side-groups. The methods of Bicerano [2] and Hopfinger [8] and for a number of situations that of Van Krevelen [i], do permit calculation of the Tg-value of polymers with different kind of side- groups. The effect of 'long' paraffinic side-chains, however, can easily incorporated in the 'modified cohesive energy' method. It is wellknown that these paraffinic side-chains decrease the Tg-value and it is possible to calculate the Tg- values of these systems using a proper side-chain length depending reduction factor (the same method is followed by Van Krevelen). Vinyl polymers with linear side-chains are represented by:

-[CH2-~]m) where X stands for a trivial structural unit; (C~H~ n the polymer with n = 0 is called the

basic polymer.

The following series of Tg-values can be derived with some effort from the graph reported by Eisenberg [13]:

alpha- poly-alkyl polymethyl olefins styrenes acrylates Tg,~ Tg,~ Tg,~

n = 0: -21 I00 105 n = I: -31 78 65 n = 2: -46 - 35 n = 3: -56 6 20 n = 4: -61 - - n = 5: -66 -26 -5

n = 7: - -44 -22 n = 8: - -52 - n = 9: - -65 - n = Ii: - - -65

The ~Ecoh. value of the -CH- and -CH3 groups from the -CH(CH2- CH3) side-chain of poly(1-butene), is according to Table 7.1: 3430 + 9826 = 13256 J/mol. This value has to be multiplied with 0.896 to decrease the calculated Tg-value from -21~ for polypropylene to -31~ for poly(1-butene), (n = 1). The reduction factors calculated in this way are plotted as a function of the number of CH2 units in Figure 7.1. The third order polynominal curve fit has a correlation coefficient of 0.9844 and permits calculation of the following average reduction factors:

n = 1, 0.86 n = 5, 0.47 n = 2, 0.74 n = 6, 0.41 n = 3, 0.63 n = 7," 0.36 n = 4, 0.54 n = 8, 0.32

Values for n higher than 8 are meaningless; in chapter 1.4.2 it is shown that the Tg-value determination for samples with n > 7 is seriously hampered by side-chain crystallisation effects.

236

Ef fec t of the number of CH2 units in a polymers' s i de -cha in on the T g - v a l u e

Figure 7.1

1 . 0 0

0.90

0 .80

0 .70

L

0 0 .60 0

c 0 .50 0

o u -

O 0 .40 "0

rr 0 . 3 0

0.20

0 .10

0.00

- +

_ + +

+

0 , . I . , I , I , I , I , -

2 4 6 8 10 12

Number of CH2 units in the s ide-cha in

237

The Tg-value of poly(l-heptene), for example, is then calculated as follows :

The repeating unit is : - [CH2-CH ( {CH2 }4-CH3) ] n-

The -CH(CH3)- unit contribution is reduced by 0.54 (n = 4). The mainchain consits of sequential C-atoms, hence the n z 5 ~Ecoh. values from Table 7.1 have to be used:

~Ecoh. -CH- ~Ecoh. -CH3

~Ecoh. -CH2-

: 3430 J/mol. : 9826 J/tool.

13256 J/mol. x 0.54 = 7158 J/mol. : 4936 J/mol.

Ecoh. = 12094 J/mol.

En, i - 2, using equation 7.1 results then in-

Tg, calc. = 0.0145 x (12094/2) + 120 = 208 K

A Tg-value of 212 K is reported in literature, see Table 7.2.

The Tg-value of poly(ether ether ketone), PEEK is finally calculated as a second example of a Tg-value calculation according to the 'modified cohesion energy' method-

The repeating unit is : - [O-C6H4-O-C6H4-CO-C6H4] n-

The mainchain does not contain sequential C-atoms, hence the n _< 4 ~Ecoh. values from Table 7.1 have to be used:

~Ecoh. -C6H4-: 27951 hence 3 x 27951 = 83853 J/mol. ~Ecoh. -CO- : 29243 hence 1 x 29243 - 29243 J/mol. ~Ecoh. -0- : 4852 hence 2 x 4852 = 9704 J/mol.

Ecoh. = 122800 J/tool.

En, i = 6, using equation 7.1 results in:

Tg, calc. = 0.0145 x (122800/6) + 120 - 417 K

The measured value was also 417 K, see Table 7.2.

The Tg-values of 48 polymers were calculated, see Table 7.2, using the 'modified cohesion energy' method as shown above. The Tg-values of twenty-seven of these systems were used to calculate ~Ecoh. values of the different structural groups i.e. these values are fitting by definition. The average (Tgexp. - Tgcalc.) value of the other (21) independently calculated Tg-values is • 8 K. This value is greater than • 8 K (between 9 and 30 K) for only six of these systems.

Table 7.2 also lists the Tg-values calculated according to Van Krevelen's method and Bicerano's method. Thus, a direct comparison of the results of these three methods is possible. The figures 7.2, 7.3 and 7.4 show the three Tg(experimental)/ Tg(calculated) correlations of the twenty-one (for the 'modified cohesion energy' method independently calculated)

238

Table 7.2 Results of Tg-value calculations Tgc, 1: group contribution method (Van Krevelen) Tgc, 2 : connectivity indices method (Bicerano) Tgc,3: 'modified cohesion energy' method (*)

polymer name/structure Tgc, 1 Tgc, 2 Tgc,3

K K K

Tg-value experim.

K

�9 polybutadiene, cis trans

- [CH2-CH=CH-CH2 ] n-

170 172 164 237 - 181

164 181

(a) (a)

2. polyoxymethylene, - [CH2-O] n-

223 215 191 191 (b)

3. polyethylene, - [CH2-CH2 ] n-

195 187 192 195 (c)

4. polyoxltrimethylene, - [CH2-CH2-CH2-0] n-

209 201 191 195 (b)

5. polyisobutylene, - [CH2-C (CH3) 2 ] n-

196 190 203 203 (a)

6. polyisobutyleneoxide - [ CH2- C (CH3) 2-0] n-

199 204 (a)

7. polyisoprene, cis - [ CH2 - (CH3) C=CH-CH2 ] n-

196 202 206 205 (a)

8. C2C3 alternating copolymer - [ CH2 - CH2 - CH2 - CH ( CH3 ) ] n-

230 215 222 211 (a)

9. poly (l-heptene) - [CH2-CH ({C~2 }4-CH3) ] n-

213 190 208 212 (d)

10. poly (1-hexene) - [cs2-c. ( { c~2 } 3-cH3) ] n-

218 196 216 217 (d)

11. polyisoprene, trans 218 220 (C)

12. poly (cis-chloroprene) - [ CH2-CH=C (C1) -CH2 ] n-

233 237 218 225 (e)

13. poly (1-pentene) - [cM2-c. ( { c~2 } 2-cH3) ] n-

226 203 227 227 (d)

14. polyvinylidene fluoride - [CH2-C (F) 2]n-

206 242 233 233 (b)

15. poly (1-butene) - [CH2- CH (CH2-CH3 ) ] n-

238 216 238 236 (a)

(*) Tg-value, K = 0.0145�9 (Ecoh./En, i) + 120

239

Table 7.2 continued (2) Tgc, 1: group contribution method (Van Krevelen) Tgc, 2 : connectivity indices method (Bicerano) Tgc, 3- 'modified cohesion energy' method

polymer name/structure Tgc, 1 Tgc, 2 Tgc, 3

K K K

Tg-value experim.

K

16. polypropylene, atactic - [CH2-CH (CH3) ] n-

255 235 252 252 (a)

17. polyvinylidene chloride - [CH2-C (CI) 2]n-

255 276 253 255 (b)

18. polyethylene succinate 236 - [0-CO- (CH2) 2-C0-0- (CH2) 2]n-

254 279 272 (f)

19. poly(vinylpropionic acid) - [ CH2 - CH (0 - { CO } - CH2 - CH3 ) ] n-

295 282 275 (a)

20. poly(methyl acrylate) - [CH2-CH (CO-O-CH3) ] n-

279 308 283 283 (b)

21. poly (4-methyl pentene- 1) - [CH2 - CH (CH2 - CH { CH3 } - CH3 ) ] n-

2 4 2 2 8 0 286 ( a )

22. PP/CO alternating copolymer - [CH2-CH (CH3) -COl n-

281 292 297 2 9 4 (a)

23. poly (vinyl acetate) - [ CH2 - CH (0- CO- CH3 ) ] n-

302 307 303 303 (a)

24. poly(3-methyl butene 1) - [CH2 - CH (CH { CH3 } - CH3 ) ] n-

259 300 307 (a)

25. polyvinyl fluoride - [CH2-CH (F) In-

328 284 314 314 (b)

26. nylon 6.6 - [ (CO) NH- (CH2) 6 - N H (CO) - (CH2) 4] n -

329 323 323 (a)

27. poly (p-xylylene) - [ CH2 -CH2- (C6H4) ] n-

361 329 303 333 (e)

28. poly(ethylene terephthalate) 369 - [0 (CO) - (C6H4) - (CO)O- (CH2) 2]n-

373 340 348 (a)

29. poly (chloro-p-xylylene) - [ CH2 -CH2 - (C6H3CI) ] n-

407 350 366 353 (e)

30. poly(vinylalcohol) - [CH2-CH (OH) ] n-

357 338 353 353 (b)

31. poly (vinylchloride) - [CH2-CH (C1) n-

354 293 354 354 (a)

240

Table 7.2 continued (3) Tgc,l: group contribution method (Van Krevelen) Tgc, 2 : connectivity indices method (Bicerano) Tgc, 3 : 'modified cohesion energy' method

polymer name/structure Tgc, 1 Tgc,2 Tgc,3

K K K

Tg-value experim.

K

32. poly (oxy [p-phenylene] ) - [ (C6H4) -0] n-

393 359 358 358 (b)

33. polyformal -[R-O-CH2-0]n- R = - (C6H4) - (CH3) C (CH3) - (C6H4) -

407 415 361 361 (a)

34. poly [thio (p-phenylene) - [ (C6H4) -S]n-

347 357 363 363 (g)

35. phenoxy resin - [R-O-CH2-CH (OH) -CH2-O] n-

399 388 368 368 (a)

36. poly (acrylonitrile) - [ CH2 - CH (CmN) ] n-

364 373 373 (b)

37. polystyrene - [CH2-CH (C6H5) ]n-

373 379 377 378 (a)

38. styrene/CO altern, copolymer 362 - [CH2-CH (C6H5) -CO] n-

381 380 383 (a)

39. poly (methyl methacrylate) 378 - [CH2- (CH3) C (CO-O-CH3) ] n- atac.

357 340 384 (h)

40. poly (ether ether ketone) 420 - [0- (C6H4) -0- (C6H4) -CO- (C6H4) ] n-

433 417 417 (a)

41. polycarbonate - [R-O- (CO) -0] n-

409 419 420 423 (a)

42. Udel - [R-O- (C6H4) -S02- (C6H4) -0] n-

465 462 469 458 (e)

43. dichloro polycarbonate PC with one - (C6H2C12)- unit

475 459 459 (i)

44. tetramethyl polycarbonate, PC 461 with two - (C6H2 [CH3] 2)- units

494 498 476 (i)

45. poly(dimethyl PPO) - [ (c6.2 { cs3}2 ) -o] n-

46. tetrachloro polycarbonate PC with two -(C6H2C12)- units

386 437 483

509 498

483 (j)

493 (i)

241

Table 7.2 continued (4) Tgc, 1 : group contribution method (Van Krevelen) Tgc, 2 : connectivity indices method (Bicerano) Tgc, 3 : 'modified cohesion energy' method

polymer name/structure Tgc, 1 Tgc, 2 Tgc, 3 Tg-value experim.

K K K K

47. poly (arylene sulphone) - [ (C6H4) -0- (C6H4) - S02 ] n-

469 483 493 493 (e)

48. tetrabromo polycarbonate, PC - with two - (C6H2 [Br] 2) - units

511 523 523 (i)

(a) Tg-values determined according to the procedure described in 1.1.3.

(b) Tg-values reported in the Polymer Handbook [14] (c) The Tg-value of PE is still under discussion. We accept

Van Krevelen's value of 195 K [i]. (d) see ref. [13]. (e) see ref. [2]. (f) see ref. [5]. (g) see ref. [15]. (h) Bosma et.al, report a Tg-value of 383 K for atactic PMMA

with a Mn value of 45000 [16]. Min and Paul [17] report a Tg-value of 386 K (both DSC onset values measured at a heating rate of 20~ Hence, we are using a value of 384 K.

(i) see ref. [18]. (j) see ref. [19].

continued from page 237: Tg-values. The correlation coefficient is 0.9740 for the group contribution results, 0.9835 for the (Bicerano) connectivity indices results and 0.9948 for the 'modified cohesion energy' results. The results of the 'simple' modified cohesion energy method are for linear polymeric systems clearly improved in comparison with the other two methods. It is important to realise, however, that the group contribution method of Van Krevelen and especially the connectivity indices method of Bicerano are covering clearly more structural possibilities.

Figure

242

Tg(cal.)/Tg(exp.) correlation (group contribution method) 7.2

5 1 0

4 7 0

4 3 0

~" 3 9 0

> 3 5 0 I

O " ~ 3 1 0

2 7 0

2 3 0

190 L 190

C o r r e l a t i o n c o e f f . - 0 . 9 7 4 0

+

+

+ / +

++/+

+

2 3 0 2 7 0 3 1 0 3 5 0 3 9 0 4 3 0 4 7 0 510

Tg(exp.)-value, K

Tg(

cal.

)-va

lue,

K

= BO

�

9 W

=

0 0

LO

+,

, 0 I'0

0 ~o

-,,,i

0 GJ

cD

-~

x 0

I < 0

......

c CO

0 -I~

0 -,,1

0

I'0

CO

CO

W

-I~

-I~

"q

= O1

�

9 ~

"-,1

0 0

0 0

0 0

' '

I'

''

I "

' I

''

" I

' i

' I

'

_ +.

....

+ o

- +

"-~.

8

" II

0

0 C -7 CD

A 0 o -I

CD

N

0 --

<--I

CD

D_ ~--

0 Co

0

3~

d)

B}

D"~"

0

2 4 4

Tg(cal.)/Tg(exp.) (modified cohesion

Figure 7.4

correlation energy method)

5 1 0

4 7 0 Correlation ooe f f . = 0 . 9 9 4 8

4 3 0

3 9 0

> 3 5 0 I . _ , . . .

(3 v

3 1 0 +

+

2 7 0

2 3 0

19O 190 2:30 2 7 0 3 1 0 350 3 9 0 4 3 0 4 7 0 510

Tg(exp.)-value, K

245

7.2.3 The T_u-value of cross1~nked polymeric systems Crosslinking reduces the molecular mobility and thus causes increased Tg-values. This increase of the Tg-value due to crosslinking is often described in literature by the semi- empirical equation [5, 20, 21, 22]-

Tg = Tg, o + k/Mc 7.2

where: Tg = the Tg-value of the crosslinked system, Tg, o = the Tg-value of the non-crosslinked system with

a Mn value z 50000. k = a constant and Mc = the molecular weight in between the crosslinks.

For crosslinked elastomers i.e. rubbers vulcanised with sulphur, Mc can be written as:

MC = Mn/p = (Mn.f)/(2.v') 7.3

where: ~ = the crosslink density, f = the network functionality v' = the network chain density

Hence, 7.2 can be written as :

Tg - Tg, o = (2 .v, .k) / (f.Mn) = C.v' 7.4

Thus, the Tg-value difference increases according to equation 7.4 linearly with the network chain density. This relation was checked for two vulcanised rubber systems; an experimental solution SBR (SSBR) and an commercial emulsion SBR (ESBR, Into1 1502). The experimental values measured for the SSBR system are listed in Table 7.3.

Table 7.3 Tg-value increase of a SSBR system due to vulcanisation with sulphur (gum vulcanisate samples).

- . , , , - T ,.,lq

sulphur concentration

. . . . . . . . . ,,,

0 . 0 , , , , , , , ,

1 . 0

1.25 ,, .,.

' 1.50 L

2.0

3.0 , ,

6.0 . . . . . . . . ,

10.0

Tg-value, Tg-Tg, o

aC ~ mol./m3

-40.0 0.0 .,, ,

-37.0 3.0 , ,

-36.5 3.5 , . , .

-35.5 4 .S u , i ,

-34.0 6.0

-32.5 7.5 , ,,,.,,

-25.0 15.0 . . . . . . . i

-15.0 25.0 ' " I . I I ' I ' . . . . . ' , , 'P I

V' -value,

,, .,

53.3 _

76.7 i _ .,. __1

96.7 , ,,,, , , , L J . ,

131.7

181.7 _

340.0

580.0 ! , j - ! ,! , ! ~ _

246

The Tg increase of SSBR due to vulcanisation versus the network chain density

Figure 7.5

25

20

~ 1 5 (!) (..) E

. . . . .

r

-~10 > I E~ F-

5

0 , . . I ~ I I , I . . , I I I ,

0 100 2 0 0 :300 4 0 0 5 0 0 6 0 0

Del ta T g - 0 . 0 4 3 v ' + 0.4 cor re la t ion c o e f f . - 0 . 9 9 9 3

Network chain density, mol./m3

247

The Tg-value increase plotted as a function of the network chain density in Figure 7.5, shows indeed a linear relation which nearly crosses the origin for the non-crosslinked, pure rubber. Hence, equation 7.2 decribes the Tg-value increase of an experimental SSBR system due to vulcanisation satisfactorily.

The Tg-value of the emulsion polymerised SBR sample Into1 S1502 increases from -57~ for the pure, non-vulcanised polymer to -33"C after vulcanisation with 10.0 phr of sulphur. The experimental values measured for this system fit considerably less good with the linear relation predicted by equation 7.4 than the values measured for the SSBR system:

ESBR ~Tg = 0.035 x v' + 2.26, correlation coeff. -- 0.9827

SSBR ~Tg = 0.043 x v' + 0.30, correlation coeff. -- 0.9993

The reason for this difference in behaviour of these two rubbers might be the difference in molecular structure. Mainly linear rubber molecules are polymerised in the solution process, while the emulsion process produces a product with a considerable amount of branched polymer chains.

It seems that equation 7.2 already fails to describe adequatly the Tg-value increase due to crosslinking of high molecular weight linear polymers if branching disturbs the regularity of the formed network. It is then not realistic to expect that this relation will describe the far more complex, three- dimensional chain building process of low molecular weight resins and their curing agents.

The diglycidyl ether of bisphenol A (DGEBA) which at room temperature is highly viscous but still liquid, is used in many epoxy resin applications. A three-dimensional network is formed after reaction of DGEBA with a proper three or four functional curing agent. The structure of p/LE~DGEBA is given below:

CH3 R I H2 C-CH-CH2-0 O - C - ~ O- CH2 -HC~-~H2

CII3

The purity of this 'basic building block' is indicated by its functionality i.e 2.00 and by its epoxy molar mass (EMM) value i.e. 340/2 = 170.

The commercially available DGEBA products (EPIKOTE 828, Shell Chemicals - DER 332, Dow Chemicals or MY/GY 250 - Ciba-Geigy) are not completely pure DGEBA systems and their structure is, therefore, more realistic given by-

248

,o ~H ,o, H2C-~H-CH2- [O-R-O-CH2-CH-CH2 ] n-O-R-O-CH2-HC-CH2

CH3

where n is about 0.1 and R represents- - -

CH3

The purity of the DGEBAused influences the regularity of the network formed after the cure reaction and this is reflected in Tg-value differences. This effect is clearly shown by a series of experiments in which an epoxy resin was purified in a number of steps. These purified DGEBA resins were cured, subsequently, with a stoichiometric amount of DDM as curing agent:

H2N~-CH2~-NH2, 4-4' diaminodiphenylmethane (DDM)

The DSC onset Tg-values (see 1.1.3) of the complete, carefully cured systems are listed below.

Table 7.4 Effect of the impurity level of DGEBA resins on the properties of the resin and on the Tg-value of the cured systems.

resin type

rw~m purity, % EMM-value functionality

Tg-value, ~ Tra-value, ~ Hf-value, J/g

, , . ,

DGEBA/DDM Tg-value, ~

A

91.9 185 1.8

154

92.1 184.5 1.971

-17 39

0.1

168

C

93.8 181.2 1.984

-21 48 72

171

D

99.6 170.6 2.000

-22 46 85

186

A: EPIKOTE 828 EL commercial epoxy grade, B: experimental (improved purity) EPIKOTE 828, C: purified version of B, D- purified version of C. systems cured: 1 hour/80eC - 1 hour/150~ - 1 hour/175~ - 0.5 hour/200aC, the weight losses due to the 200eC cure-step were ~ 2.0 %wt.

249

The resins A and B are liquid resins at room temperature while the resins Cand D are highly crystalline systems under the same conditions. The strong influence of purity i.e. the EMM value, and the functionality on the Tg-value of the (DDM) cured end-product is clear. The Tg-values of cured end- products are also influenced by differences in the resin/ curing agent mixing procedure and by differences in the curing schedule. All these effects make it difficult to speak about 'the' Tg-value of a certain resin/curing agent system and hamper a proper Tg-value estimation.

7.2.4 The Tg-value estimation of epoxy_ resin based. crosslinkedresin systems The network obtained after the cure of the DGEBA with DDM (each DDM molecule reacts with four DGEBA molecules) can be represented by:

/X -N ~m OH 9H

CH2 - CH - CH2 - 0 - R - 0 - CH2 - CH - CH2~ N ~ - CH2 ~ N~ CH2 - CH - CH2 - Y

In such a network alternating DDM/DGEBA molecule chains can be recognised. Hence, the network can considered to contain long, linear molecule chains with as 'repeating unit':

X

- [CH2- -CH2-O-R-O-CH2-CH-CH2- -CH2 -N] n-

X

Each of these 'repeating units' is connected on two places via the crosslinks with the other molecule chains. The Tg-value of this system can considered to be the sum of the Tg-value of a 'linear polymer' with the above given repeating unit and a crosslinking effect. The Tg-value contribution can be calculated using the 'modified cohesion energy' method with the n ~ 4 ~Ecoh. values from Table 7.1:

~Ecoh. -CH2- : 4936 hence 5 x 4936 = 24680 J/mol. ~Ecoh. -CH- : 3430 hence 2 x 3430 = 6860 J/mol. ~Ecoh. -OH : 28708 hence 2 x 28708 = 57416 J/mol. 6Ecoh. -0- : 4852 hence 2 x 4852 = 9704 J/mol. ~Ecoh. -R- : 85007 hence 1 x 85007 = 85007 J/mol. 6Ecoh. -C6H4-: 27951 hence 2 x 27951 = 55902 J/mol. ~Ecoh. -N- : 4183 hence 2 x 4183 = 8366 J/mol.

Ecoh. = 247935 J/mol.

~n,i = 16, using equation 7.1 results in: Tg-value , 0.0145 x (247935/16) + 120 = 345 K

250

The highest Tg-value measured for a DGEBA/DDM system is 459 K (see Table 7.4). Hence, the Tg-value contribution of two crosslinks in the above given 'repeating unit' is 459 - 345 = 114 K i.e. the maximum Tg-value contribution per crosslink is 57 K.

Resin system D was, subsequently, cured with 4-4' diamino diphenylsulphone (DDS), 4-4' diaminodiphenylpropane (DDP) and 4-4' diaminophenyl (DP) to check the consistency of the ' crosslink contribution' .

resin D/DDS system, ' repeating unit' : X

- [CH2-CH-CH2-O-R-0-CH2-CH-CH2- - - ] n-

X Tg(exp.) - 499 K Tg(cal.) -- 383 K i.e. (499 - 383)/2 = 58 K per crosslink

resin D/DDP system, 'repeating unit': X

_ - n -

X Tg(exp.) - 469 K Tg(cal.) - 367 K i.e. (469 - 367)/2 = 51 K per crosslink

resin D/DP system, 'repeating unit': X

- [CH2 CH2-O-R-O-CH2- I X

Tg(exp.) = 452 K Tg(cal.) - 343 K i.e. (452 - 343)/2 = 55 K per crosslink

Hence, the average Tg-value contribution per crosslink is 55 • 3 K for resin system D. The Tg-value of resin system A/DDM is 427 K. The Tg-value contribution per crosslink is for this system (427 - 345)/2 ~ 41 K/crosslink. Hence, if an 'epoxy resin'/DDM system is used to determine the Tg-value contribution per crosslink, the Tg-values of this 'epoxy resin' with other curing agents can be calculated. In other words, the correction for the resin impurity and preparation effects on the Tg-value of the end-product occurs via the Tg- value contribution per crosslink.

251

The network obtained after the reaction of resin system A/HHPA (hexahydrophthalic anhydride) can be represented by the structure given below.

X I ~:o

9

o o-~,, ~ ~-o- -~,-c~,-o-~-o-~,-~-o- ~-o-~,

H2 \ / \ /

0-. q:o ~ - - g-o

~ 80 O-R-O-~,~--~ -o-'~ ~-o- - 0 - CH2 - CH- CH2 - 0 - - - CH- CH2 - - CH2 \ / ~H2

C> * x_. / C-O I

X

The ' repeating unit', containing two crosslinks, is considered to be:

X I

- [ c H - o - ~ ~ - o - c H 2 - c ~ - c ~ 2 - O - R - O - C H 2 ] ~ -

x 0 The calculated Tg-value of a linear polymer with this repeating unit is 321 K. The contribution of two crosslinks is 82 K. Hence, the calculated Tg-value for the resin system A/ HHPA system is 321 K + 82 K = 403 K. The measured Tg-value was 399 K. The Tg-values of five other epoxy resin systems were calculated (and measured) in the same way, all results are listed in Table 7.5. It is important to reaiise that the used structures are 'idealised' structures. Every deviation from this ideal network structure will lower the experimental Tg- value. Thus, the calculated Tg-values will in general be equal or higher than the measured Tg-values. This was confirmed for the six systems investigated.

252

resin system

Table 7.5 Results of Tg-value calculations on crosslinked epoxy resin systems

Tg-vaiue .... number of Tg-value rep. unit, crosslinks/ calc.,

K rep. unit K �9 ,,, , ,,, .,,

Reference resin A/DDM 345 6Tg/crosslink = 41 K 427

, . , ,,

resin A/E 112 ~ 341 1 382 365 332 2 414

resin A/E 113 353 2 435 432

resin A/HHPA 321 2 403 399 resin A/TMA 317 3 440 435

resin A/BABA 346 i 3 469 445 resin A/ 346 1 387 377 Alnovol

, , ,, L, ,, ,

* two seemingly nearly equal possibilities

Tg-value meas.,

K .,

The structural formulas of the systems in Table 7.5 are given below.

EPIKURE 112 ~-NH- CH2 -CH2- CH2-NH2

EPIKURE 113 H2N-O -CH2- O-NH2

CH3 CH3

BABA

\CH2 ~ - NH2

0 HHPA Q ~ ~\0 and TMA

Alnovol VPN 1981 =oQ c==Q o= / ,crt2

HO~CH2-(~ -OH

=~176 G )o

253

7.3 The Tm-value estimation

7.3.1 Introduction The application of polymers in all kind of end-use products is often limited by their relative low melting temperatures. There has been therefore a considerable interest in determining the factors which control the Tm-value. One would hence also expect an extended amount of Tm-value/chemical structure correlations proposed in literature, just as the many correlations proposed for the Tg-value, see 7.2. It is remarkable however, that this is not the case. Van Krevelen [1] seems to be the only one with a clear Tm-value/chemical structure correlation concept. Even Bicerano [2], who is reporting such an extended polymer properties prediction system, does not mention the Tm-value at all.

The Tm-value of a polymeric system is according to Young [23] and Cowie [4], depending on: i. mainchain symmetry, ) 2. mainchain flexibility, ) molecular structure 3. intermolecular bonding/polarity, ) effects 4. tacticity and cis/trans isomerism, ) 5. type and size of side-groups, ) 6. molar mass, 7. thermal history (crystal size/perfection), 8. the presence of residual monomer/solvents, 9. rate and type of measurement.

A number of these Tm-value depending factors are also mentioned in the earlier mentioned list concerning the Tg- value (see 7.2.1); in fact there is a close agreement between the two lists. Hence, some relation between Tg and Tm can be expected and is indeed reported (the Tg/Tm ratio) in literature [24]. This Tg/Tm ratio appears to vary widely however, but indicative values might be:

Tg/Tm - 0.5 , for symmetrical polymers, Tg/Tm - 0.67, for unsymmetrical polymers.

These values can only be used as a 'rule of thumb' not as a real correlation. Van Krevelen argued that this correspondence suggests that a treatment analogous to that proposed by him for the Tg-value could also be used for the prediction of Tm- value. Subsequently, he calculated the group contributions (Ym) resulting in a system to estimate Tm-values of polymers with widely varying molecular structures.

254

7.3.2 The reduced Tu/Tm correlations One of the problems hampering the Tm-value/chemical structure correlations is in the experimental reference values. Most experimental values reported are non-equilibrium results while equilibrium Tm-values, the so-called Tm(o)-values, are preferred. Van Krevelen does not mention the difference between equilibrium and non-equilibriumTm-values and used probably mainly non-equilibriumTm-values, simply because equilibrium Tm(o)-values for many polymers are not (yet) available.

A series of non-equilibrium Tm-values were measured on samples recrystallised from the melt under standard conditions (the so-called Tm2-values, see 1.1.4) and a number of non- equilibrium literature values were used to look for an improved correlation between Tm/Tg relations. We tried to improve the results of such a relation by distinguishing different groups of polymers instead of looking for one relation for all types of polymers. Three groups of polymers offering the best fitting correlations, were selected finally:

Group A, polymers with cis/trans double bonds, vinyl polymers with side groups/chains and polymers with p-phenyl groups.

Group B, polymers with strong intermolecular bonding effects. Group C, linear polymers with - (C) n- or - (C) n-0- mainchains.

The Tm/Tg relations for these three groups of polymers (see Figure 7.6) appeared to be practically linear. The straight lines calculated have nearly identical slopes and differ only in the constant value added. Thus, it is possible to approximate the Tm/Tg relation for these 29 polymer systems satisfactorily by one relation with different constant values:

Tin-value (K) = 1.23 x (Tg-value, K) + To 7.5

where To - 74 for group A polymers, 152 for group B polymers and, 206 for group C polymers.

The calculated Tm-values with this 'reduced Tm/Tg correlation' method, the Tm-values calculated with Van Krevelen's group contribution method and the experimental Tm-values are listed in Table 7.6. The figures 7.7 and 7.8 show the Tm(exp.)/ Tm(calc.) correlations for both methods. The correlation coefficients of these two correlations and the average [Tm(exp.)-Tm(calc.)] values are:

Tm (exp.) -Tin (calc.), average value

group contribution reduced Tm/Tg method (v. Krevelen) correlation

38 K _+ 35 K 15 K+II K

Correlation coefficient 0. 819 0. 981

255

Tg/Tm relation

6 2 0

5 8 0

5 4 0

5 0 0

r 4 6 0 :::3

> I 4 2 0 E

I--

3 8 0

18 polj4n~rs

Figure 7 .6

+

0

Z~

Gro~ B 0 6 Dolj4ners

+

+

+

+

+

Gro~ C 5 Dol~ners

+

3 4 0 ~

3 0 0 +

2 6 0 ......... J ......... ,

A: Tm = B: T m - C: T m -

1 . . . . . . . / .....

1 . 2 3 x T g + 1 . 1 5 x T g + 1.1 l x T g +

7 4 1 6 9 2 3 2

1 6 0 2 0 0 2 4 0 2 8 0 3 2 0 3 6 0 4 0 0 4 4 0

Tg-value, K

256

Table 7.6 Results of Tm-value calculations Tmc, 1 : group contribution method (van Krevelen) Tmc,2" reduced Tm/Tg relations method, equation 7.5

polymer name/structure Tg- Tmc, 1 Tmc, 2 Tm-value value experim.

K K K K

Group A _Dol~rs; Tmc,2 = 1.23xTg + 74

1. polybutadiene, cis trans

- [ CH2-CH-CH-CH2 ] n-

164 179

359 276 295

265 (a) 315 (a)

2. polyisoprene, cis - [ CH2- (CH3) C=CH-CH2 ] n-

205 315 326 300 (b)

3. polyisobutylene, - [CH2- C (CH3) 2 ] n-

203 316 324 318 (c)

4. poly (1-pentene) - [CH2- CH ( { CH2 } 2- CH3 ) ] n-

227 311 353 348 (c)

5. poly (1-butene) - [CH2- CH (CH2-CH3 ) ] n-

236 361 364 397 (a)

6. poly (cis-chloroprene) - [Ch2-CH=C (Cl) -CH2 ] n-

225 377 351 353 (c)

7. polymethylmethacrylate (i) - [CH2- (CH3) C (CO0-CH3) ] n-

329 (c) 473 479 442 (d)

8. poly(butene terephphalate) 307 - [O-CO- (C6H4) -CO-O- (CH2) 4] n-

488 452 473 (e)

9. poly (oxy [p-phenylene] ) - [ (C6H4) -0] n-

358 559 514 4 9 0 (e)

i0. poly(vinylalcohol) - [CH2-CH (OH) ] n-

353 539 508 503 ( f )

11. poly(ethylene terephthalate) 348 - [O-C0- (C6H4) -C0-O- (CH2) 2]n-

528 502 516 (a)

12. polyformal - [R-0-CH2-0]n- R - - (C6H4) - (CH3) C (CH3) - (C6H4) -

361 656 518 524 (a)

13. polystyrene (s) - [CH2-CH (C6H5) ] n-

378 516 539 532 (a)

14. poly(styrene/CO) copolymer - [CH2-CH (C6H5) -CO] n-

383 498 545 554 (a)

15. poly(thio[p-phenylene] ) - [ (C6H4)-S] n-

363 559 520 548 (e)

257

Table 7.6 continued (2) Tmc, I- group contribution method (Van Krevelen) Tmc, 2 : reduced Tm/Tg relations method, equation 7.5

Tg- Tmc, 1 Tmc, 2 Tin- value value experim.

K K K K

16. polycarbonate - JR-O-CO-O] n-

423 659 594 563 (e)

17. poly(ether ether ketone) 417 - [0- (C6H4) -0- (C6H4) -CO- (C6H4) ] n-

576 587 617 (a)

Group B polymers: Tmc,2 = 1.23xTg + 152

1. polyvinylidene fluoride - [CH2-C (F) 2]n-

233 437 439 443 (C)

2. polyvinylidene chloride - [CH2-C (CI) 2]n-

255 462 466 473 (c)

3. polyvinyl fluoride - [CH2-CH (F) ] n-

314 530 538 503 (c)

4. nylon 6.6 323 - [ (CO) NH- (CH2) 6-NH (CO) - (CH2) 4] n-

540 549 534 (a)

5. polyvinyl chloride - [ CH2 -CH (CI) ] n-

354 576 587 583 (C)

6. poly (acrylonitrile) - [CH2-CH (C=N) ] n-

373 598 611 614 (c)

Grou_D C _Dolors: Tmc, 2 = 1.23xTg + 206

1. polyoxy (isobutylene) - [CH2- (CH3) C (CH3) -0] n-

204 435 457 450 (C)

2. polyoxymethylene 191 426 441 451 (a) - [CH2-O] n-

3. polyformaldehyde - [CH (CH3) -0] n-

200 602 452 452 (g)

4. polypivalolactone 260 - [CH2- (CH3) C (CH3) - (CO) -0] n-

433 526 511 (a)

5. ethylene/C0 copolymer 257 418 522 528 (a)

258

Table 7.6 continued (3)

(a) Tm2-values, determined according to the procedure described in I.I.4.

(b) see ref. [I]. (c) Tin-values reported in the Polymer Handbook [14]. (d) see ref. [17]. (e) series of equilibrium Tin-values reported by Cheng and

Wunderlich [25]. They report for PEEK a Tin-value of 668 K, we measured a Tm-value of 617 K. Hence a difference of 51 K. They report for PET a Tm-value of 553 K, we measured a Tm- value of 516 K. Hence a difference of 39 K. The difference between the equilibrium and the non-equilibrium Tm-values seems to be in the order of 45 K for these types of polymers. The experimental Tin-values for PPO, PBT, PO, PPS and PC are the values of Cheng and Wunderlich corrected with 45 K.

(f) see ref. 26. (g) see ref. 24.

259

Tm(ca l . ) /Tm(exp . ) c o r r e l a t i o n (group c o n t r i b u t i o n me thod )

620 Figure 7.7

580

540

500

-~ 460 > I

-~ 420 0 E F- 380

340

3OO

260

+

+

~ +

Rval + + +

+

I ,, I

+ +

+

+

+

++ + +

260 300 340

+

- - . - - . - -

, I , I , , , I , I I

380 420 460 500 540

0 . 8 1 9

1 580 6 2 0

Tm(exp.)-value, K

260

Tm(cal.)/Tm(exp.) correlation (reduced Tm/Tg relations method)

Figure 7.8

6 2 0

580 +

r :D

. - - . - - ,

> I

O v

E I--

540

500

460

420

380

340

3.00

+

+

+

+

+

+

R v a l . =

+ +

+

+

+

0 . 9 8 0 8

260 260 3 0 0

, I I , , , I _ f I ,

3 4 0 380 4 2 0 4 6 0 500 I I

540 580 620

Tm(exp.)-value, K

261

The correlation between Tm(exp.) and Tm(calc.) and thus the average difference between Tm(exp.) and Tm(calc. ) are indeed improved. The average difference is however nearly twice the average [Tg(exp.) -Tg(calc.) ] value, see 7.2.2.

Hence, this 'reduced Tm/Tg correlation' system is far from ideal. Besides, each of the three groups of polymers contain a number of strongly deviating systems. The isotactic vinyl polymers (group A) which form helical structures for instance deviate strongly. The [Tm(exp.)-Tm(calc.)] values of a number of this type of polymers are:

poly(1-butene) : 32 K, Tin(exp.) = 397 K. poly(1-propene) : 44 K, Tin(exp.) = 429 K. poly(3-methyl-l-butene) : 119 K, Tin(exp.) = 576 K. poly(4-methyl-l-pentene) : 62 K, Tin(exp.) = 504 K.

Poly(oxy[di-methyl p-phenylene]), PPO is such a strongly deviating system for the high Tg, high Tm-value polymers. The Tg-value of 483 K results according to equation 7.5 in a Tm(calc. ) -value of 668 K. The experimental Tin-value is, however, 535 K i.e. a difference of 133 K.

Polytetrafluoro ethylene is such an exception for the group B polymers. Cheng [27] reports for PTFE a Tg-value of 200 K and a Tm-value of 605 K. Equation 7.5 results in a calculated value of 398 K. PTFE behaves perhaps, by its molecular symmetry, like a group C polymer; in that case the calculated Tm-value should be 452 K. Hence, there remains a difference of at least 153 K.

The Tm-value of polyethylene poses a problem for the group C polymers. The Tg-value of 195 K results in a Tm(calc.)-value of 446 K while we measured an experimental value of 405 K for HDPE. However, the Tg-value of PE is still under discussion. Alberola et al. [28] recently suggested again a Tg-value for PE of about 158 K. The corresponding Tm-value calculated of 400 K agrees reasonably with the experimental value.

The Tm-values of polyethylene oxide and polypropylene oxide are also difficult to calculate. The determination and calculation of the right Tg-values of these systems seems to be the main problem.

The difference between the equilibrium Tm(o)-values and the non-equilibrium Tin-values was mentioned already. The lack of sufficient reliable Tm(o)-values hampers the determination of proper Tg/Tm(o) correlations. For the polymers of group A however a number of Tm (o) -values are available [25, 27]. These data are plotted in Figure 7.9 as a function of the corresponding Tg-values and result in the following relation-

Tm(o)-value (K) = 1.37 x (Tg-value, K) + 71 7.6

The ten calculated and experimental Tm(o)-values are listed in Table 7.7. The average [Tmo(exp.)-Tmo(calc.)] value is 23 K _+

262

10 K. Hence, a proper set of 'reduced Tm/Tg correlations' seems to offer better fitting calculated Tm values than the more general method proposed by van Krevelen. The differences between the calculated and the experimental Tmvalues are for both methods still to high. Thus, more and more reliable Tg- and Tm(o)-values are necessary to develop better Tm(o)/ chemical structure correlations.

Table 7.7 Results of Tm(o)-value calculations.

polymer name/structure Tg- Tin(o) - Tin(o) - value value value

calc. exp. K K K

Group A _Dol_vmers: Tm (o), c = 1.37xTg +71

1. polybutadiene, cis - [ CH2-CHffiCH-CH2 ] n-

164 296

2. polyisobutylene - [CH2- (CH3) C (CH3) ] n-

203 349

3. poly (1-pentene) - [ C H 2 - C H ( { CH2 } 2 - C H 3 ) ] n -

2 2 7 382

4. poly (1-butene) - [CH2-CH (CH2-CH3 ) ] n-

236 394

5. poly(butene terephthalate) 307 - [O-CO- (C6H4) -C0-O- (CH2) 4 ] n-

492

6. poly(oxy[p-phenylene] ) - [ (C6H4) -0] n-

358 561

7. poly(ethylene terephthalate) 348 - [O-CO- (C6H4) -CO-O- (CH2) 2]n-

548

8. poly (thlo [p-phenylene] - - [ (C6H4)-Sin-

363 568

9. polycarbonate - [R-0-C0-0]n- R= - (C6H4) - (CH3) C (CH3) - (C6H4) -

423 651

10. poly(ether ether ketone) 417 - [0- (C6H4) -0- (C6H4) -CO- (C6H4) ] n-

642

2 8 5

317

403

411

518

535

5 5 5

593

608

668

I

> I 0

v

E F-

Figure 620

580

5 4 0

500

4 6 0

4 2 0

380

3 4 0

300

2 6 0 L 1 6 0

263

Tg/Tm(o) (group A

re lat ion polymers)

7.9

-I- 4-

-I-

+

-I- + ,

+ Tm(o) - 1 . 3 7 x T g + 71 R v a l . - 0 . 9 7 8

1 9 5 2 3 0 2 6 5 3 0 0 3 3 5 3 7 0 4 0 5 4 4 0

T g - v a l u e , K

264

7.4 The Hf-value estimation

Important changes in the physical properties of polymers are observed during passage of the glass-rubber transition region and/or the fusion region. The temperature location of these two regions is indicated by respectively the Tg-value and the Tm-value. The level of most physical polymer properties, especially in the temperature region between Tg and Tm, is mainly determined by the extent of the crystalline phase (the crystallinity).

The crystallinity or the fraction of crystalline material x(c) can be derived from specific volume data, from specific heat data, from infrared extinction coefficient data, from X-ray scattering data, from NMR data or from heat of fusion data. The heat of fusion of a polymeric system as determined by DSC can give an x(c)-value"

x(c) - (Hf)/Hf(max.) 7.7

where Hf(max.) is the heat of fusion value of the completely crystallised system. Since polymers cannot be completely crystalline, Hf(max.) has to be determined by extrapolation. Reliable Hf(max.) values are scarce and this hampers the crystallinity determination by DSC. However, Hf-values as such can give an indication of the crystallinity level. Besides, this determination is frequently used to compare different samples or to compare identical samples after different thermal treatments.

The maximum attainable degree of crystallisation during spontaneous, spherulitic crystallisation of flexible polymers under quasi-isotropic conditions depends to a great extent on the maximum rate of crystallisation Iv(max.)]. The v(max.) is related to the Tg/Tm ratio and this relation permits Van Krevelen [1] to report a x(c) versus Tg/Tm curve. Van Krevelen also reports a purely empirical expression for v(max.) based on the observation that the growth rate is high if the regularity of the molecular structure is strong:

v(max.) = 83x[(n.CH2/Z) . ( I / { l + I~ , } ) ] ' 7.8

where: n - the number of CH2 or equivalent groups in the backbone of the structural unit,

Z = number of atoms in the backbone of the structural unit,

- number of carbon atoms in the side group.

Because of lacking Hf(max.)-values, the Hf(1)-values (see 1.1.4) of fifteen polymers as such were correlated with both their Tg/Tm-values and with their v(max.)-values calculated according to equation 7.8. The relevant data are listed in Table 7.8 and plotted in Figure 7.10. The line drawn in the Hf-value/Tg/Tm-value plot can only suggest an indicative 'maximum possible' Hf-value considering the scattering data points.

265

Figure 300

270

240

210

180 c~

150 4-: I

The

7 . 1 0

Hf -va lue versus crystal l isat ion

A Ny/c,,~ O PET 6.6

the maximum rate 4" PPL A B R

120

90

60

3O

0" 0.001

H I - ~ I ~ ~ l h ~ T ~ I ' ~

4 l#

2 0 0 ' - + ' '

1 8 0 4

1 4 0

" 1 2 0 o4h

- ~ l O 0

O0 _ L

6 o +§ o +

- 0 . 4 0 0 . 5 0 0 . 6 0 0 . 7 0 0 . 8 0

. T g / T m

" § A

0.01 0.1 1 10 100

v(max.), /Jmeter/s

266

Table 7.8 Tg/Tm-, v(max.)- and Hf (1) -values of a series of polymers

polymer name/structure Tg- 'I'm- Tg/Tm v (max.) Hf (I) value value value value value

K K ~m/s J/g

1. polyethylene - [CH2-CH2 ] n-

195 405 0.48 83 195

2. polyoxymethylene - [CH2-O] n-

191 451 0.42 83 184

3. polypropylene - [CH2-CH (CH3) ] n-

252 436 0.58 0.32 95

4. poly (1-butene) - [CH2-CH (CH2-CH3 ) ] n-

236 397 0.59 0.064 60

5. poly(eth, eth. ketone) 417 - [ { O- (C6H4) } 2 -CO- (C6H4) ] n-

617 0.68 0.026 57

6. poly (3-methyl butene- 1) 307 - [CH2-CH (CH{G"H3 }2} ]n-

576 0.53 0.020 48

7. poly (trans 2.3 epoxy- 242 butane - [CH (CH3) -CH (CH3) -O] n-

420 0.58 0.013 59

8. poly (styrene/CO) - [CH2-CH (C6H5) -CO] n-

383 554 0.69 0.013 50

. poly (4-methyl 286 pentene- 1 ) - [ CH2 -CH ( CH2 -CH { CH3 } 2 ) ] n-

504 0.57 0.008 36

10. polystyrene (s) - [ CH2 -CH (C6H5) ] n-

378 532 0.71 0.004 22

11. polyformal - [R-O-CH2-O]n- R = - (C6H4) -C (CH3) 2 - ( C6H4 ) -

361 524 0.69 0.004 10

12. poly ethylene 348 terephthalate (PET) - [0-C0-(C6H4) -CO-0- (CH2)2]n-

516 0.67 2.1 60

13 .Nylon 6.6 323 534 - [ CO - NH- ( CH2 ) 6 - NH - CO- ( CH2 ) 4 ] n -

0.60 21.6 73

14 .polypivalolactone (PPL) 260 - [ CH2 - C (CH3) 2 - CO- O] n-

511 0.51 5.2 103

15.polybutadiene (BR, cis)164 - [ CH2 -CH=CH-CH2 ] n-

265 0.62 0.61 33

267

Table 7.8 continued

* -0- = -CH2- * -CO- = -CH2- (only for styrene/CO copolymer) * -(C6H4)-, Z -- 4 * -(C6H5) , E = 5 * -C(CH3) 2, E = 2 conjugated between two aromatic rings * -C(CH3) 2, E - 0 non-conjugated, aliphatic chains.

Now many v(max. )/Hf-value data points seem to give, however, a reasonable correlation (Rval. is 0.9907). This correlation is described by the equation-

Hf-value (J/g) ~ 38.Sxlog v(max.) + 115.5 7.9

This equation suggests that too low Hf-values are measured for Nylon 6.6, PET, PPL and BR. This indicates that this correlation is giving a kind of 'maximum possible' Hf-value and that the experimental values will be equal or smaller than the values calculated.

268

7.5 The thermal stability estimation

7.5.1 IntrQduction The thermal stability of a polymer, that is the potential of its chemical bonds to withstand thermal energy, depends in general strongly on the test conditions. Usually, mainly four different situations occur: - relative short-time, high temperature isothermal or non-

isothermal conditions, - relative long-time, low temperature isothermal conditions

and, - both above conditions performed in an inert and/or in an

oxidative (air) atmosphere. Relative long-time experiments are often performed according to the Underwriters Laboratories (UL) testing protocol [29]. The results, the retention of certain properties upon exposure to heat, are expressed in UL temperature index values. The time and material consuming UL testing procedure is usually preceded by short-time mass retention measurements, mainly performed with TGA equipment.

Thermal stability/chemical structure correlations as reported by Van Krevelen [1] and Bicerano [2] are using and predicting the relative short-time, thermal stability values measured in an inert atmosphere. The mass losses measured under these (non-isothermal) conditions are caused by the damage of chemical bonds due to chain depolymerisation and/or random decomposition (see 2.2.1). Van Krevelen [1] distinguishes five experimental indices which characterise this non-isothermal decomposition process : - the initial decomposition temperature (Td, o), at which the

loss of weight is just detectable, - the temperature of half decomposition (Td,1/2), at which the

loss of weight reaches 50% of its final value, - the temperature of the maximum rate of decomposition

(Td, max. ), - the average energy of activation (Eact. ,d), determined from

the temperature dependence of the rate of weight loss, - the amount of char residue at the end of the experiment

(usually at a standard temperature of 900~

Van Krevelen and Bicerano obviously consider the Td,1/2-value as the most important, for they both report a Td,1/2 - chemical structure correlation. Van Krevelen mentions however, that these decomposition indices are interrelated i.e. Td,o is about 0.gxTd, 1/2, Td,max. is about Td, 1/2 and Eact.,d is about Td,1/2 - 423.

We prefer to use the Td,o value as a measure for the short- time thermal stability of a polymer. This because-

- the properties of a polymer sample that has lost 50% of its mass differ considerably with that of the initial product and,

269

- the Td, o - Tm temperature region offers a more realistic value for the 'processing window' of a polymer than the Td, 1/2 - Tm region.

Detection of the exact temperature at which the ma~s loss process starts is experimentally far more difficult than determination of a Td, i/2-value. Besides, the determination of the Td,o-value might be hampered by mass losses due to residual solvent/ monomer and/or oligomers. We expected that a kind of compromise would be possible by using semi-statically determined Td, o-values as shown already for polypropylene in chapter 2.

7.5.2 The semi-static Td.o-value determination In chapter 2.2.2 it was shown that the isothermally determined (static) Td, o-value of polypropylene (PP) is about 190~ This value is approached reasonably in a dynamic i.e. non- isothermal experiment if a heating rate ~ 0.1~ is applied. A clear separation is obtained during such an ultra- low heating rate experiment between the mass losses due to evaporation of the oligomers fraction and mass losses due to the thermal decomposition process(es). The step-wise change in the mass/temperature curve of Figure 2.6 at relative low temperature is considered to be due to evaporation of the oligomers fraction (the amount observed agreed with the values determined isothermally). The temperature or temperature region at which the zero mass loss rate or a very low, constant mass loss rate changes into clearly increasing mass loss rates with increasing temperatures, is considered to be the beginning of the overall thermal decomposition process (the Td, o-value). In connection with the ultra-low heating rate used (0.1"C/minute) this Td, o-value is called the semi- static Td.o-value.

Figure 7.11 shows the result of a semi-static Td, o-value determination of poly(4 methyl-l-pentene). Mass losses & 0.05 %wt. can easily be recognised; in line with the TGA balance resolution and sensitivity. The weight loss rate is practically zero between 30~ and II0~ Between about II0~ and 267~ a small, but nearly constant weight loss rate is detected (0.003 %wt./oc). These mass losses might be caused by the evaporation of a small oligomers fraction (0.5 %wt. maximally). The mass loss rate increases strongly at temperatures > 267oc due to what is considered to be the overall thermal decomposition of the polymer sample.

7.5.3 Thermal stability estimation based on Td.o-values Van Krevelen [I] reports a correlation between the chemical structure and the Td, i/2-values of polymers:

Td, I/2 (K) = (ENi x Yd, I/2.i)/M 7.10

which is called the molar thermal decomposition function.

SamDle Weight' 4.5BI Po ly- (4methy Ipenteen-1)

mg

I01. O0

I00 50

100 O0

99.50 -I

99.00

J~ 98.50

98. O0

bO ,,,3 O

97.50

97. O0 Figure 7.1 1

96.50 The Td,o determination

96.00 50.0

nitrogen atmosphere TEMPI" 30~ TIME1: O.Omin TEMP2: 500"C

tO0.0

RATE1: 0.1~

i50.0 200.0

Temperature (~

250.0

267~

300.0

271

Table 7.9

polymer name/structure

1. polyethylene - [ CH2- CH2 ] n-

Td,o-values and Yd, o-values of a series of polymers

Td, o- group Yd, o- Yd, 1/2 - value value value*

K K.kg/mol K.kg/mol

592 -CH2- 8.3 9.5

2. polypropylene - [ CH2 -CH ( CH3 ) ] n-

463 -CH(CH3) - Ii. 2 18.5

3. polystyrene, atactic - [CH2-CH(C6H5) In-, s

504 -CH (C6H5) - 44.1 56.5 538 -CH (C6H5) - 47.7

4. polyisobutylene - [ CH2 - C (CH3) 2 ] n-

496 -C (CH3) 2- 19.5 25.5

5. polybutadiene 1.4, cis 542 - [ CH2-CH=CH-CH2 ] n-

- CH=CH- 12.7 18

6. polyisoprene, cis 487 - [ CH2 - C ( CH3 ) =CH- CH2 ] n -

-C (CH3) =CH- 16.5 21.5

7. poloxymethylene 415 - [CH2-O] n-

-0- 4.2 8

8. poly(pivalo lactone) 516 - [CH2-C (CH3) 2-C0-O] n-

-CO- 19.7 14

9. poly ethylene tere- 532 -CO- (C6H4) -CO- 77.2 phthalate

- [O-CO- (C6H4) -CO-O- (CH2) 2]n-

103

10. poly butylene tere- 544-CO-(C6H4)-CO- 78.2 phthalate

- [O-CO- (C6H4) -CO-O- (CH2) 4 ] n-

103

11. polycarbonate - [O-R-O-COIn-, R = - (C6H4) -C ( CH3 ) 2 - (C6H4) -

527 -R- 105.8 143

12. poly (l-butene) - [ CH2-CH (CH2- CH3 ) ] n-

542 -CH(CH2-CH3) - 22.1

13. poly (3-methyl- 1-butene) - t cH2 -cH (c~ { cH3 } 2 ~ 3 n-

546 -CH3 7.9

14. poly(4-methyl- 540 1-pentene) -Ec~2-c~ (CH2-CH(C.3 }2~ J.-

-CH3 7.5

* values reported by Van Krevelen [1], Table 21.2 p. 646.

272

van Krevelen also indicates that there is a relation between Td,1/2 and Td, o (see 7.5.1). We therefore decided to look for the correlation between Yd, o group contribution values based on our semi-static Td, o values and the Yd,1/2 group contribution values reported by Van Krevelen. The semi-static Td,o-values of fourteen polymers were determined, and used to calculate semi-static Yd, o group contribution values. These semi-static Yd, o-values were then correlated with the Yd,1/2- values reported by Van Krevelen, see Figure 7.12. The correlation proved to be fairly good, with its correlation coefficient of 0.9953. ore reasonable to assume that (for a first estimation) Yd, o group contribution values can be calculated using Van Krevelens' Yd,1/2 group contribution values with the equation:

Yd, o = 0.74 x Yd, 1/2 + 1.5 7.11

The Td, o-value of poly(ethylene terephthalate), for example, is then calculated as follows:

repeating unit PET, - [O-CO- (C6H4)-C0-0- (CH2) 2In- and M s 192

-CO-(C6H4)-C0- - 77.7 K.kg/mol i.e. 1 x 77.7 = 77.7 K.kg/mol -(CH2)- = 8.3 K.kg/mol i.e. 2 x 8.3 - 16.6 K.kg/mol -0- - 4.2 K.kg/mol i.e. 2 x 4.2 = 8.4 K.kg/mol

102.7 K.kg/mol

Td, o-value, calculated = (102.7 x 1000)/192 - 535 K The experimental Td, o-value is 532 K. The Td,o-value of 535 K is 185 K lower than the calculated Td,1/2-value (720 K, [i]).

The Td, o-value of poly (2,6-dlmethyl p-phenyleneoxide) is calculated subsequently �9

repeating unit PPO, -[(C6H2{CH3}2)-O]n- and M-- 120 g/tool

Yd, i/2 -(C6H2{CH3}2)- - 82.0 K.kg/mol [i], Yd,o is then according to equation 7.11:

Yd, o = 0.74 x 82 + 1.5 = 62.2 K.kg/mol Yd, o -0- = 4.2 K.kg/mol

66.4 K.kg/mol

Td,o-value, calculated = (66.4 x 1000)/120 = 553 K. This Td, o- value is 197 K lower than the calculated Td,1/2-value (750 K, [ :z)) .

The Td, o-values of the most common polymeric systems can thus be calculated using equation 7.11 and Van Krevelens Yd,1/2- values.

u

0 E

v_

D

>

c 0 0

D 0 k_

I o

>-

120

100

80

60

40

20

0

Semi-s ta t ic

_ +

I _ I I

Yd,o

4-

. I .

- Y d , 1 / 2

Yd,o

......... L .... I .... n ! +

correlation

- 0 .74xYd,1 /2 + R v a l . - 0 . 9953

1.5

Figure 7.12 Yd, 1/2-group contr.

0 20 40 60 80 100 120 140 160

values, K.kg/mol

,L,,, I o I _ I I

274

7.6 The moisture sensitivity estimation

Low molecular weight components such as residual solvents and/or monomers have a strong influence on the physical properties of a polymeric system. The presence of water forms a special case in this respect. Firstly because nearly every polymer absorbs a certain amount of moisture from the air. Secondly because of the special character of the water molecule, i.e. relative small but with a strong tendency towards hydrogen bond formation in its own liquid as well as with other polar groups.

The equilibrium water saturation of a polymeric system increases with the number of polar groups present in the polymeric matrix. Circumstances like the accessibility of the polar groups, the relative strength of the water-water versus water-polymer bonds and for semi-crystalline polymers the degree of crystallinity, hamper a straightforward correlation between the number of polar groups and the solubility. Van Krevelen [1] presents the amount of water per structural group at equilibrium (expressed as molar ratio), as what he calls the best possible approach to the sorptive capacity of (amorphous) polymers versus water.

For many polymers it takes a long time to reach a real equilibrium water saturation. Resin samples, for instance, need immersion times of at least eighty days to reach a real equilibrium water saturation conditions, see chapter 5.2. The immersion times necessary to reach equilibrium water saturations increase to more than threehundred days for glass fibre filled, resin laminate systems. Based on these relative long immersion times it was expected that also for nonpolar polymers like polyethylene very long immersion times are necessary to reach real equilibrium conditions. Hence, we determined a number of structural group molar water content values by a series of long-term water absorption experiments.

Nine sample strips (about 100x10x0.1 mm) of the polymers listed in Table 7.10 were first dried at 50~ in vacuum until constant sample weights were reached. The dry samples were stored subsequently for more than two years in ion-free water (in the dark) at 20 • 1~ Then the weight increase due to water absorption was measured on surface-dry samples. The samples and the equilibrium water saturation values measured are listed in Table 7.10.

The molar water content of these polymers per structural group was calculated using these equilibrium water content values. The contribution of the -R- and -(C6H4)- groups was assumed to be zero to calculate the contributions of respectively the -O- CO-O- group and the -S02- group. Hence, these contributions will be slightly too high. The molar water content values per structural group calculated are listed in Table 7.11 together with the values published by Van Krevelen [1].

275

Table 7.10 The equilibrium water saturation values measured after more than two years of water immersion.

polymer name/structure crystalline equilibrium molar water phase water content

estimated saturation, calculated % wt. % wt. for group

1. Alkathene-2 LDPE - [ CH2 -CH2 ] n-

37 0. 0047 -CH2 -

2. Carina 531 PVC - [CH2-CH (C1) In-

0 0.0543 -CH (Cl) -

3. KM6100 polypropylene - [CH2-CH (CH3) ] n-

48 0. 0772 -CH (CH3) -

4. Dow 666 polystyrene - [ CH2 -CH (C6H5) ] n-

0 0. 0721 -CH(C6H5) -

5. Penton

- [ CH2 -C (CH2CI) 2 -CH2 -0] n- 0 0.1633 -O-

6. poly(pivalo lactone) - [CH2-C (CH3) 2-C0-0] n-

60 0. 3460 -CO-

7. Lexan GE polycarbonate - JR-O-CO-O] n- R = - (C6H4) -C (CH3) 2 - (C6H4) -

0 0. 3922 -O-C0-0-

8. polysulphone 0 - [O- (C6H4) -S02- (C6H4) -O-R] n-

0.7431 -S02-

9. Akulon nylon 6 - [ (CH2) 5-NH-CO] n-

50 5. 8515 -NH-CO-

a. Hf-value LDPE : 103 J/g, Hf (max.) PE : 276 J/g i.e. x(c) -- 0.37

b. Hf-value PP : 90 J/g, Hf(max.) PP: 188 J/g i.e. x(c) - 0.48

c. The crystallinity values of PPL and Nylon 6 are estimated values.

276

Table 7.11 Maximum molar water content values per structural group at 20~

I S t ructual group

-CH2-

I -CI -CH (CH3) -

I -CH(C6H5)- ..

-0-

I -c 0- - o - c o - o - .

I ~. 'S02-. , ..

measured molar water content value

5.8E-5

1.8E-3

3.4E-3

4.1E-3 .. ..

1.0E-2 ,,,

3.8E-2 , ..,

5.5E-2

1.5E-1

7.3E-I Ill I II

I I

molar water i content value Van Krevelen

.., ....... , .

(5.0B-s) I . ,,

(~.0E-I) I ,

(I.0E-4) 1 5.0E-3 _ I 1.0E-I

3.0E-I I , . , , .,

i - i

2.0E0 I II I

The values measured for the -CH2- group and the -CH(C6H5)- group agree reasonably with the values of Van Krevelen. All the other new values are however clearly lower.

The maximum equilibrium water saturation of a cured epoxy system (resin system A/HHPA, see 7.2.3) was calculated succesfully using these new molar water content values. As ' repeating unit' was recognised for this three dimensional network structure (see 7.2.4) :

- [CH-O-CO- (C6HI0) -CO-O-CH2-CH-CH2 -O-R-O-CH2 ] n-

where R = -(C6H4)-C(CH3)-(C6H4)-

Using the measured molar water content values from Table 7.11 results in:

2 x -CO- : 2x3.SE-2 tool. water/per repeating unit 4 x -O- : 4x1.0E-2 mol. water/per repeating unit 5 x -CH2/CH- �9 5x5.8E-5 mol. water/per repeating unit 3 x -C6H4/C6H10-: 3x4.1E-3 tool. water/per repeating unit

0.1286 mol. water/per repeating unit i.e. 0.1286x18 = 2.32 gram of water per 464 gram of resin. Thus 0.50 gram of water per 100 gram of resin or 0.50 %wt.

The equilibrium water saturation measured of this epoxy resin casting sample was 0.90 %wt. Using the molar water content values reported by Van Krevelen ([1], Table 18.11 page 572), an equilibrium water concentration of 2.4 %wt. is calculated.

277

This result seems to confirm these new, lower experimental molar water content values. But it will be clear that also the results calculated with this small number of new molar water content values can be only indicative ones. More experimental results like shown in Table 7.10 are necessary to obtain really firm molar water content group data.

7.7 Estimation of the key-properties of a new polymer

Chapter nine shows how the different TA techniques focussed on one product, contribute to the characterisation of a new polymeric system- aliphatic polyketone. This polyketone is a perfectly alternating copolymer from carbon monoxide and ethylene (PK copolymer). The key-properties of this new polymeric system were calculated as an example.

Repeating unit : - [CH2-CH2-CO] n-

1. The Tg-v~lue

The mainchain consists of sequential C-atoms, hence the n _> 5 6Ecoh. values from Table 7.1 have to be used.

~Ecoh. -CH2- 2x 4936 = 9872 J/tool. 6Ecoh. -CO- lx18405 = 18405 J/mol.

Ecoh. = 28277 J/mol.

En, i = 3, using equation 7.1 results in.

Tg(calc.), K = 0.0145 x (28277/3) + 120 = 257 K

2, The Tm-v~Aue

The linear mainchain makes PK copolymer a group C polymer. The Tin-value is then according to equation 7.5:

Tm(calc.), K - (1.23 x 257) + 206 = 522 K

3 J The Hf-value

The v(max.) value is, according to equation 7.8:

v(max.) -- 83 x (2/3) 4 = 16.4

The Hf-value is then, according to equation 7.9, equal or smaller than:

Hf-value, J/g _< 38.5 x log(16.4) + 115.5 = 162 J/g

278

4. The Td. o-value

The Yd, o value of PK copolymer is (2 x 8.3) + 19.7 = 36.3 K.kg/mol. According to equation 7.10 the Td, o temperature is then calculated as-

Td, o-value - (36.3 x 1000)/56 = 648 K

5. The maximum water saturatiop

Using the new molar water content values from Table 7.11 results in:

2 x -CH2- : 2xS.8E-5 mol. water/repeating unit 1 x -CO- : 1x3.8E-2 mol. water/repeating unit

0.03812 mol. water per repeating unit i.e. 0.03812 x 18 - 0.686 gram of water per 56 gram of polymer.

Thus 1.23 gram of water per 100 gram of completely amorphous PK copolymer can be expected. Hence, the estimated maximum water saturation of the semi-crystalline PK copolymer as such (assuming that only the amorphous phase absorbs water) is about 0.5 %wt.

6. Comparison of the calculated and measur%d values

A computer program was developed to perform the calculations necessary to obtain these key-propertles. A print-out of the results obtained with this program is shown in Figure 7.13. The values calculated above are shown again, but the program also performs a data base search and gives the experimental values, if available for comparison (values between brackets).

The calculated Tg-value is the Tg-value of the completely amorphous polymer. Thus, the experimental Tg-value has to be equal or higher than the estimated value due to the presence of a considerable crystalline phase (PK copolymer, x(c) = 0.63, see 9.2.2). The agreement between the estimated and measured Tg values seems therefore not too bad.

The agreement between both estimated crystalline phase values (the Tm- and Hf-value), and the experimentally obtained values is also not too bad.

The estimated Td, o-temperature giving an indication of the polymers' thermal stability, holds only for polymers with the more or less 'simple' thermal degradation processes of chain depolymerisation and/or random decomposition. S. Shkolnik and E.D. Well [30] reported, however, that for non-stabilised PK copolymer a process of furan ring formation can start at temperatures of about 250 ~ Water, one of the reaction products, is coming free during this intramolecular

279

cyclisation process. This reaction water will evaporate during a TGA scan and will be detected as a mass loss effect. Hence, the results of the TGA semi-static Td, o-value determination and that of the Td, o-value calculation can not be compared in this case.

The calculated equilibrium water saturation of 1.23 %wt. holds for completely amorphous PK copolymer i.e. if x(c) = 0.63, the equilibrium water saturation value of the semi-crystalline polymer as such, becomes 0.50 %wt. The equilibrium water saturation, measured on sheet material in demineralised water at 20~ was 2.55 %wt. The calculated water saturation depends in this case mainly on the molar water content value of the CO group used. The calculated value is about five times too low and this illustrates again the necessity of looking for better defined molar water content values.

These results show that the existing chemical structure/ physical properties relations still need to be considerably improved to become real development ' tools '. Improved measuring/calculating techniques (like the modulated DSC and computer modelling) and measurements on well defined series of polymers might result in a clear improvement of these chemical structure/physical properties relations.

FIGURE 7.13 Print-out of the results of the polymer key- p r o p e r t i e s c a l c u l a t i o n p r o g r a m

*************************************************************************

*** EXSYS4, P o l y m e r K e y - p r o p e r t i e s e s t i m a t i o n ***

Repeating unit : -[ (CH2)2-CO-]n

Polymer name : PK copolymer

Calculated Tg-value

Estimated Tm-value Estimated Hf-value i.e. a semi-crystalline polymer

Estimated Td, o-value i.e. processing window

Estimated Eq. water saturation (estimated EWS of the amorphous phase only)

: -16 C (4)

: 249 C (258) < 162 J/g (152)

: 375 C ( - ) : 127 C

: 1,225 %wt ( - )

( ) reference data

280

References

I. D.W. van Krevelen" Properties of Polymers, third edition, Elsevier, Amsterdam, 1990.

2. J. Bicerano: Prediction of Polymer Properties, Marcel Dekker Inc., New York, 1993.

3. J.T. Seitz, J. of Appl. Pol. ~c., Vol. 49, (1993), p. 1331 - 1351.

4. J.M.G. Cowie: Polymers: Chemistry & Physics of Modern Materials, Int. Textbook Company Ltd., Aylesbury, Bucks., 1973.

5. U.T. Kreibich and H. Batzer, Die Angew. Makromol. Chemie, 83, (1979), p. 57 - 112.

6. R.F. Fedors, Pol. Eng. & Sc., 14, (1974), p. 147 - 154. 7. D.R. Wiff and M.S. Altieri, J. of Pol. Sc.: Polymer

Physics Edition, Vol. 23, (1985), p. 1165- 1176. 8. A.J. Hopfinger et al., J. of Pol. Sc. : Polymer Physics

Edition, Vol. 26, (1988), p. 2007 - 2028. 9. V. Bellenger et al., J. of Pol. Sc.: Polymer Physics

Edition, Vol. 25, (1987), p. 1219 - 1234. 10. Yong-GU Won et al., J. of Pol. Sc. : Polymer Physics

Edition, Vol. 29, (1991), p. 981- 987. 11. Xinya Lu and Bingzheng Jiang, Polymer, Vol. 32, 3, (1991),

p. 471 - 478. 12. R.F. Boyer, Macromolecules, 25, (1992), p. 5326 - 5330. 13. A. Eisenberg, J.E. Mark and W.W. Graessley: Physical

properties of Polymers, American Chemical Society, Washington (DC), 1984.

14. J. Brandrup and E.H. Immergut: Polymer Handbook, Wiley, New York, third edition 1989.

15. S.Z.D. Cheng, Zong Quan Wu and B. Wunderlich, Macro- molecules, 20, (1987), p. 2802- 2810.

16. M. Bosma et al., Macromolecules, 21, (1988), p. 1465. 17. K.E. Min and D.R. Paul, J. of Pol. Sc., Part B: Polymer

Physics, 26, (1988), p. 1021- 1033. 18. A.F. Yee and S.A Smith, Macromolecules, 14, (1981), p. 54

- 64. 19. S.Z.D. Cheng and B. Wunderlich, Macromolecules, 20,

(1987), p. 1630- 1637. 20. T.G. Fox and S. Loshaek, J. Polym. Sci., 15, (1955),

p.371. 21. F. Rietsch, Polymer, 17, (1976), p. 859. 22. L. Banks and B. Ellis, Polymer, 23, (1982), p. 1466. 23. R.J. Young: Introduction to Polymers, Chapman and Hall,

London (1983) . 24. R.F. Boyer, Rubber Chem. Techn., 36, (1963), p. 1303. 25. S.Z.D. Cheng and B. Wunderlich, Thermochimica Acta, i/~,

(1988), p.161- 166. 26. Y. Nishio and R. St. J. Manley, Macromolecules, 21,

(1988), p. 1270 - 1277. 27. S.Z.D. Cheng: Polymer Analysis and Characterisation,

Editor H.G. Barth, J. Wiley, New York (1989). 28. N. Alberola el al., Eur. Polym. J., 28, (1992), p. 935 -

948.

281

29. Modern Plastics Encyclopedia, McGraw-Hill, New York (1989) p. 659.

30. S. Shkolnik and E.D. Weil, Journal of Applied Polymer Science, Vol. 69, (1998), p. 1691- 1704.

TG-VALUES OF POLYMERS WITH DOUBLE BONDS IN THE

MAIN CHAIN

CHAPTER 8

282

CHAPTER 8: Tg-VALUES OF POLYMERS WITH DOUBLE BONDS IN THE MAIN CHAIN AND Tg-VALUES OF NON-POLAR POLYMERS WITH SIDE-CHAINS

8.1 Introduction

The presence of double bonds in the mainchain of a polymer increases the number of appearance forms of such a polymer. Well known examples of such systems are polybutadiene rubber (BR) and polyisoprene rubber (IR). BR for instance, can be polymerised into the following configurations:

CH2 CH2 [/ \ / \]n

CH--CH

CH2 / \] n CH2

CH-CH [/ \ /] n [\ / CH CH2 I

CH=CH2

1,4 cis-BR 1,4 trans-BR 1,2-BR

But, 1,2- or vinyl BR can be polymerised in the atactic, the syndiotactic or the isotactic form. Hence, five different configurations can be obtained by polymerisation reactions with butadiene (CH2=CH-CH=CH2) as monomer. The product obtained depends on the catalyst system used but is usually a mixture of 1,4 cis-, 1,4 trans- and atactic 1,2-BR. The commercial processes using Co-, Ni- or Ti-based catalyst systems, for instance, produce BR with a 1,4 cis-BR content higher than 90 %wt. But butyllithium initiated homopolymerisation of butadiene results in a product with 1,4 trans-BR contents up to 60 %wt.

All these commercially produced BR systems are amorphous rubbers under atmospheric conditions. The Tg-value of these systems, depending on their structure, is described by the Gordon-Taylor relation, see Chapter 1. BR becomes a semi- crystalline polymer under atmospheric conditions if the 1,4 trans-BR content is higher than about 70 %wt. or if a syndio- or isotactic 1,2-BR phase is present. This is shown by the results of thermo-analytical measurements on experimental BR systems with a high trans content and with a high syndiotactic 1,2-BR content which are reported in this chapter. Moreover, the Tg-values of two series of IR samples containing 1,2- and 3,4-IR are used to determine the Tg/structure relation for non-polar polymers with side-chains.

8.2 Experimental BR systems

8.2.1 BR with a hiuh 1.4 trans content Five BR samples with 1,4 trans-BR contents between 60 %wt. and 90 %wt. (as determined by FTIR) were prepared with an anionic, Ba containing catalyst system. These systems wer~ heated in the DSC from 20~ to 80~ to detect the presence of a possible

283

crystalline phase due to spontaneous crystallisation during storage under atmospheric conditions. Subsequently, the samples were cooled to -120~ and reheated to 80~ at a rate of 20~ The results of these measurements are collected in Table 8.1.

Table 8.1 Results of DSC measurements on BR with a high content of 1,4 trans-BR

sample number /ProPerty

1,4 trans, %wt. 1,4 cis, %wt.

....... vinyl, %wt.

scan I, amorph. or semi-cryst. Hf-va!u e, J/g

scan 2, Tc-value, oC

. Hc-value, J/g

scan 3, Tg-value, ~ Te-value, oC

, , ,, ,,, 'I

. ,,,

A

90.0 6.5 3.5

S . C .

- 45

37 72

-91 78 I

I

B

85.0 11.6 3.4

S . C .

- 26

25 22

-92 65

II

C

81.0 14.5 4.5 ..

S . C .

- 6

7 23

-93 57 I .

78.0 17.7 4.3

S . C .

- 1

-7 15 J _

-93 43

I' r

rill I I '

E

65.5 28.2 6.3

i

a.

-96

I ,

These data show that spontaneous crystallisation during storage under atmospheric conditions occurs for 1,4 trans-BR contents higher than about 75%wt. The amorphous character of system E does not mean that a BR system with an 1,4 trans content of 65.5 %wt. stays amorphous under all circumstances. Comparison of the (scan 1) Hf-value with the (scan 2) Hc-value of the systems C and D shows that cooling to lower temperatures promotes the crystallisation process. A decrease of the cooling rate from 20"C/minute to 10"C/minute was already sufficient to obtain also a crystalline phase in system E, see below.

Due to the shift of the baseline of the fusion endotherm the heat of fusion value of the samples 'as received' can only be indicated approximately. Figure 8.1 shows the recystallisation curves of the systems A/D. The three minima in the exothermic effect of sample A and the two minima in that of the systems B/D indicate the presence of a complex crystallisation/fusion process. The fusion process is for this reason characterised by its Te-value instead of a Tm-value. Figure 8.2 shows the fusion curves measured (at a rate of 20"C) for system E after cooling this sample at a rate of respectively 10~ 1~ and 0.1aC/minute. These three fusion curves clearly show how a lower cooling rate promotes the formation of crystallites melting at higher temperatures, i.e. larger and/or more perfect crystallites. Figure 8.3 finally, shows the strong 1,4 trans-BR content dependence of the fusion and recrystallisation temperatures.

284

3 " 0 l

2.5

2.0

A

I .S ~,, 1 r

I . O - "~ 0 "11"

0 . 5 -

0.0 I -80. 0

.; ~_ l -~LTf_ . - . - - - -~_ - f - f .

t

(81 X) (85::D (9BZ)

( ) t r O n o o o n ~ e n t

I I I "1 I I ' I I I I -50. 0 -40. 0 -30. 0 -20. 0 -10. 0 0. 0 IIX 0 20. 0 30. 0 40. 0

Figure 8.1 Temperature (~ W. de Jong Crystallization curves of high trans content BR samples

- - - 3 . 0

-2.5

.2.0

1.5

1.0

0. 5

, ~ -0.0 50, 0

iO .O

g .O

-0. 0

o o o l i n 9 e o ~ e , d o ~ C / m i n . | I II A 7 v

A

7.0 .~

e.0 ~: o

. ~ 0 ~ ID

"1"

4 . 0 , - -

3 . 0

2 . 0

1 . 0

o . o . . . . . . 7 ' , . . . . , -~oo. o -?s. o -so. o -2s. o o. o

Temperature (~ Figure 8.2 Fusion curves of BR rubber (trans content 65%) after crystallization at different cooling rates

85

Figure 8.3 The Te- and Tc- values of BR rubbers as a function of the trans-BR content

+ T e - v a l u e s A T o - v a l u e s

75

65

55

0 45

-(] 35

4-" 25 (P o E 1 5

5

- 5

- 1 5

- 2 5 +

2 8 5

j , , I ~ . . . . . . . . . . . . I I , I = i I ,

6 0 6 8 76 84 92 100

T r a n s - B R content, %wt.

286

The small, but clearly detected, increase of the Tg(onset)- value with an increasing 1,4 trans-BR content, confirms that the Tg-value of I00 %wt. 1,4 trans-BR is higher than that of I00 %wt. 1,4 cis BR. Tg-values of 1,4 cis- and 1,4 trans-BR of respectively -I09~ and -94~ obtained by extrapolation, were reported in Chapter 1.2.2. An 1,4 trans-BR content of 54 %wt. was however the highest trans content value of the results used for that calculation. The extrapolated Tg-value of 1,4 trans-BR shifts from -94"C to -92~ when taking into account these high trans content data. The extrapolated Tg-values of 1,4 cis- and 1,2-BR of respectively -I09~ and -16~ did not change significantly.

8.2.2 BR with a hiuh svndiotactic 1.2-BR content The thermal analytical properties of an experimental syndiotactic 1,2-BR were measured and compared with that of a commercial product RB 830 (ex-JSR). The results are collected in Table 8.2.

Table 8.2 Results of DSC measurements on BR with a high content of syndiotactic 1,2-BR

II ffPl I

sample .I type/pr~ scan i, fusion region, ~

Tm(1)-value , ~ . xf.(%)-value ' , J/g I

scan 2, Tc-value , ~ Hc-value , J/g

N

scan 3, Tg-value , ~ Tm(2) -value , ~ Hf (2) -value , J/g

L I II ' 'l I,I

I

experim. sample

168- 213 207 93

173 61

27* 202 59 III

RB 830 ( ex- JSR )

50- 130 96 23

63 18

-13 101 18

I

* only detectable after quenching in liquid nitrogen

The difference of more than 100~ in the Tm-values and the nearly four times higher heat of fusion indicate that the syndio-tactic 1,4-BR content of the experimental sample is considerably higher than that of the commercial sample.

The extrapolated Tg-value of 100 % atactic 1,2-BR is -16~ The measured values of -13~ and 27oc need not to be in conflict with the extrapolated value. The presence of a small and a relative strong crystalline phase for respectively sample RB 830 and the experimental sample might cause a Tg- value increase.

The mass/temperature curves of the commercial sample measured both in air and nitrogen are shown in Figure 8.4. Remarkable

287

I00. 5

I00. 0

9g. 5

99,0

0 ~: 98. 5

g8. 0

97. 5 . .

50. 0

W . u

t00.0

f ~ ' " - - ' " " ~ . AIR . . . _ . \,\

,11 :12:840~3:C C " ~ i \ ' ~

. . . . . . . . . . . . . . . . . . . . /! I I I I l I

150. 0 200. 0 250:0 300. 0 350. 0 400. 0 450. 0

Temperature (~ Figure 8.4 Mass/temperature curves of syndiotactic 1,2-polybutadiene RB830 (ex-JSR)

I00. 20

I00. 15

0.0

T1 1.860 min T2 29.900 min Y1 99.934 Wt. % Y2 100.121 Wt. %

Delta Y 0.187 Wt. % . . . . . . ; - ~ I00. I0

o ~ 100.05

~1~ I00-00

~ 99.95

'J ' . . . . . I I ' I -- 5. 0 I 0. 0 15. 0 20. 0 25. 0 30. 0

Time (minutes)

99. 90

gg. s5

gg. 80

Figure 8.5 Mass increase of a syndiotactic 1,2-polybutadiene (ex-JSR) sample during an isothermal TGA experiment at 240~ in air

288

is the small, sudden mass increase effect starting at 223~ during the measurement in air. We thought that this effect was caused by some cyclisation reaction of the BR and attempted to confirm that by additional Tg-, Tm- and Hf-value measurements.

The mass increase of an RB 830 syndiotactic 1,2-BR sample was measured as a function of time at 240~ in an air atmosphere. About thirty minutes were necessary to obtain a mass increase of 0.19 %wt. (0.22 %wt. during the non-isothermal experiment), see Figure 8.5. This sample was cooled and then placed in the DSC. A crystalline phase with a Tm- and Hf-value of respectively 91~ and 23 J/g was detected; the Tg-value of the amorphous phase proved to be -13~ Hence, the Tg-, Tm- and Hf-values were not influenced at all by the samples' heat treatment. Even the strength of the glass-rubber transition effect (0.26 J/g."C) was not changed. Thus, the detected mass increase effect can not be explained by some cyclisation reaction of the BR.

The experimental (high syndiotactic i, 2-BR content) sample was, subsequently also heated in air and in a nitrogen atmosphere in the TGA. Figure 8.6 shows the measured mass/temperature curves. This sample is first losing about 2 %wt. residual solvent but then again shows the mass increase effect during heating in air. This effect proved to be stronger than that measured for the RB 830 sample:

onset mass increase effect, ~ total mass increase , %wt. Hf-value , J/g

exp. sample RB 830 211 223

1.05 0.22 93 23

The mass increase effect correlates roughly with the heat of fusion. This seems to confirm that this effect occurs only in the syndiotactic phase of these systems. However, what really is happening with syndiotactic 1,2-BR during heating in the presence of oxygen remains still an unanswered question.

8.3 Experimental IR systems

Isoprene CH2-C(CH3)-CH=CH2 can, just like BR, be polymerised as cis 1,4-IR (natural rubber) or trans 1,4-IR. But 1,2 and 3,4 polymerisation is also possible.

CH3 - [ CH2 - ~=CH- CH2 ] n- ~H 3 -[CH2- ]n- -[CH2-QH]n-

=CH2 ~=CH2 CH3

1,4-IR, cis/trans i, 2-IR 3,4-IR

The 1,2- and 3,4-IR can in theory be polymerised in the atactic, the syndiotactic and in the isotactic form. Hence eight different IR modifications might be possible. The Tg-value of a mixture of cis 1,4-IR and (atactic) 1,2- and

100. 0

Perkin-Eimer 7 Series Thermal Analysis System IOC. 0

gg. 5 gg. 5

gg. 0

A

o~ g8.5

e-

.~ gs. 0

g?. 5

g?. 0

g5. 5

gs.o -,, 51i 0

/ /

/ !

/

AIR

\ ,

NITROGEN

\ \ \

\

.... I ......... I ........................ ~ .............................. i ........ 100.0 150. 0 200.0 250. 0

i I 300. 0 350. 0

Figure 8.6 Temperature (~ Mass/temperature curves of syndiotactic 1,2-BR-polybutadiene

gg. 0

g8. 5

g8. 0

gT. 5

g7.0

k tl g~ I - g6. o

4O0. 0 45O. 0

RATE 1" 1.0~

290

3,4-IR is described by the Gordon-Taylor relation according to:

Tg-value, K = wl.Tg(l;4) + w2.kl.Tg(l.2) § w3.k2.Tq(3.4) 8.1 wl + kl.w2 + k2.w3

where: wl - cis 1,4-IR fraction, w2 = 1,2-IR fraction, w3 = 3,4-IR fraction and

wl + w2 + w3 = 1.0

The DSC Tg(onset)-value of cis 1,4-IR is 205 K. The Tg-values of 1,2- and 3,4-IR and the constants kl and k2 were, subsequently, determined by extrapolation.

The Tg-values of a series of samples containing atactic 3,4- IR, cis 1,4-IR and only a small amount (< 3 %wt.) of 1,2-IR were extrapolated to I00 %wt. 3,4-IR to estimate the Tg-value of this material. The measured DSC Tg(onset)-values are listed bel ow:

sample 3,4-IR weight Tg-value code fraction A 0.775 273 K ( 0~ B 0.76 271 K (-2~ C 0.63 254 K (-19~ D 0.54 244 K (-29~ E(1) 0.47 235 K (-38~ E(2) 0.47 234 K (-39~

The Tg-values of these samples are plotted as a function of their 3,4-IR content in Figure 8.7 and fit a straight line described by:

Tg-value, K = 125.4.w(3,4) + 175.6 8.2

The estimated DSC Tg(onset)-value calculated by extrapolation to w(3,4) = 1.0 is 301 K. Hence, the Tg(onset)-value of 1,2-IR and the constants kl and k2 were still to estimate.

A series of six styrene/isoprene/styrene sequential blockcopolymers were analysed. The styrene content of these systems was constant (about 14 %wt.) but the cis 1,4-IR, the 3,4-IR and the 1,2-IR contents varied. The Tg-value of the IR (rubber) phases of these samples is not or hardly influenced by the polystyrene phases. Thus, these samples could be used to estimate the lacking values for kl, k2 and the Tg of 1,2- IR. The measured Tg-values of these systems are listed in Table 8.3.

The Tg-values of 1,2-IR and 3,4-IR are expected to differ not more than a few degrees. The samples 512 and D515 can therefore considered as samples with a cis 1,4-IR phase and a 3,4-IR phase. Equation 8.1 reduces then to.

3 0 2 ::s

:3

> 2 8 4 i r 03 {.....

o 2 6 6 F-

3 0 0

~ 2 8 0

2 6 0

o

I-- 240 (9 GO t:3 220

�9 oo G 2 4 8

230 I , . !

0.00

2 0 0 0

i I , a , , , A i I , '

10 20 3o 40

Wood plot for IR samples of similar 3,4/1,2 ratio 3 2 0 -

[Tg-Tg(1,4)][(1-w(c))/w(c)]

, I , J , ,

0 . 1 0

J I i / f I . i . I l . . . l , I I, I ,, =

0 .20 0.2,0

Figure 8.7 DSC Tg(onset)-value of IR as a function of the 3,4-1R content

0 . 4 0 0 . 5 0 0 . 6 0 0 . 7 0 0 . 8 0 0 . 9 0 1.00

3 ,4-1R f rac t ion

292

Tg-value, K = 205.wl + 301.k2.w3 8.3 wl + k2 .w3

where: w3 = 0.106 (D512) results in a k2 value of 0.367 w3 = 0.208 (D515) results in a k2 value of 0.443 i.e.

an average value of 0.41 will be used for k2.

Table 8.3 Composition and Tg-values of SIS systems

D512 , , ,,

D515

D513 .,

D517

D514

I II I

1,4 1,2 fraction fraction

.. ,,,

0.894 0.021 j, ,

0.792

0.596

0.348 ,

0.216

0.161

0.036

0.047 ,

0.119 ,, , , ,,,

0.137 ..

0.176 II I I

* determined by FTIR

3,4 fraction

0.085 ,, ,,,

0.172 ,, ,,

0.357

0.533

0.647 ,

0.663 fill I I

, , " I

Tg-value K

209 ,

215 , , , , , ,

228

246

263 . . .

265

The samples D512, D514, D515, D516 and D517 can considered to be a copolymer of cis 1,4-IR and a 1,2/3,4-IR copolymer for which holds- w(3,4) -4.w(1,2). Equation 8.1 can than be written as :

Tg-value, K = Tg(1.4) + [k".Tg(c) - Tu(1.4)]_ .w(c) i - [I - k"] .w(c)

8.4

where: w(c) = w2 + w3

Tg(c) = kl.Tg(l.2) + 4.k2.Tg(3.4) kl + 4 .k2

8.5

k" -- (kl +4.k2)/5 8.6

Equation 8.3 can be written as:

Tg-value, K - Tg(c) + [Tg - Tg{1.4)]. [1 - w(c)1 k" .w(c)

8.7

The plot of Tg as a function of [Tg -Tg(l,4)]. [i - w(c)]/w(c), see the figure inserted in Figure 8.7, is giving a Tg(c) value of 299.3 K and k" value of 0.415. Substitution of the known values for Tg(c), Tg(3,4), k" and k2 in the equations 8.5 and 8.6 results in Tg(1,2) = 293 K and kl = 0.44. Thus, for IR holds :

Tg(I.4)-IR = 205 K (-68~ Tg(I,2)-IR - 293 K ( 20~ kl = 0.44 Tg(3,4)-IR = 301 K ( 28"C} k2 = 0.41

293

Equation 8.1 can now be written as-

Tg-value IR, K = 205.w(I,4) + 129.w(1.2) + 123.w(3.4) w(l,4) + 0.44w(I,2) + 0.41w(3,4)

The subsequently calculated values are compared with the measured values in Table 8.4:

Table 8.4 Measured and calculated Tg-values of SIS block

I

code

D512

D515

D513

D517 ,,,

D514

D516 i

copolymers

Ir I n' ~ I

T~ (calc.)

209.4 , ,,,,, ,

214.2

225.6

246.0

261.3

269.0 I~I

llr

T~ (exp.)

209 ,,

215

228

245 ,,,,

263

265 , I

T~(c) -Tg (e)

+0.4 , ,, ,

-0.8

-2.4

+1.0 , ,,

- 1 . 7 , , , ,

+ 4 . 0 ........... M . . . . I I

The fit between the calculated and the measured Tg-values is acceptable.

8.8

8.4 A Tg/structure correlation for non-polar polymer systems with side-chains

The difference between the calculated Tg-values of 1,2-IR and 3,4-IR proved to be only eight degrees C. This relative small difference prompted us to look for a separate Tg/structure correlation to confirm this calculated difference.

The mobility of polymer mainchains is mainly determined by the barrier to rotation around the backbone carbon-carbon bonds. In polymer systems without polar groups and/or hydrogen bond effects, this barrier to rotation is primarily determined by the size of side-groups. Hence some correlation was expected to exist between the Tg-value increase due to a side-group addition and the increase in the molecular weight of the repeating unit.

The difference in Tg-value between 1,2-IR and 1,2-BR, for example, is 36 K"

1,2-IR: Tg = 293 K 1,2-BR: Tg = 257 K

This is caused by the presence of the -CH3 side-group in the 1,2-IR. Thus, a molecular weight increase of 15 gram

294

Table 8.5 Tg-value increase of linear polymers due to side- group addit ion

polymer name/structure Tg- mol. Tg-value fract. value, weight increase, mol. wt.

K rep. unit K increase

- [ CH2- CH-CH- CH2 ] n- - [ CH2 -C (CH3) =CH- CH2 ] n- cis 1,4-BR/cis 1,4-IR

163 54 205 68 42 0.28

trans 1,4-BR trans 1,4-IR

181 54 220 68 39 0.28

i, 2-Bs - [cH2-qs] n- CH=CH2

1,2-IR -[CH2-C(CH3)]n- --c~2

257 54

293 68 36 0.28

PE - [CH2-CH2 ] n- PP - [CH2-CH (CH3) ] n-

195 28 252 42 57 0.56

PE - [ CH2- CH2 ] n- 3,4-IR -[CH2-~H]n-

~-cH2 CH3

195 28 301 68 106 1.52

PE - [CH2 -CH2 ] n-

3-methyl- [CH2-~H]3 n- but ene - 1 CH3

195 28 307 70 112 1.59

PE - [CH2- CH2 ] n- PS - [CH2-CH (C6H5) ] n-

195 28 378 104 183 2.85

PS - [CH2-CH2] n- 378 poly-alpha-methylstyrene 409

104 118 31 0.15

************************

PE - [CH2-CH2 ] n- i, 2 - BR - [ CH2 - ~H] n-

CH-CH2

195 257

28 54 62 1.00

PE - [CH2-CH2 ] n- PB-I - [CH2-CH] n-

t.2-cH3

195 236

28 56 41 1.07

PE - [ CH2-CH2 ] n- 195 227

28 70 32 1.59

295

influences a repeating unit molecular weight of 53 grams i.e. a repeating unit fractional molecular weight increase of 15/53 = 0.28 is resulting in a Tg-value increase of 36 K.

The Tg-value differences and the repeating unit fractional molecular weight increases were determined in this way for eight different systems. The results are listed in Table 8.5 and plotted in Figure 8.8. A linear relation was found for these non-polar polymeric systems:

Tg increase = 55.7 x (fractional mol.wt, increase) + 23.4 8.9

The fit between with the calculated Tg-values using equation 8.9 and the experimental values is three degrees or better. The Tg-value differences of 1,2-BR, polybutene-1 and polypentene-1 with polyethylene are also plotted in Figure 8.8. in chapter 7.2 it was discussed already that linear, alifatic side-chains are acting as plasticisers, causing a Tg- value decrease in stead of a Tg-value increase. The three separate data points in Figure 8.8 clearly illustrate that the plasticising action of a linear, alifatic side-chain starts already for a -CH=CH2 side-chain (17 K too low) and a -CH2-CH3 sidechain (42 K too low respectively) and increases with increasing side-group length (-CH2-CH2-CH3, 80 K too low).

296

Figure 8.8 Tg-value increase of linear polymers due to side-group addition

/X 1 , 2 - B R 0 P B - 1 4" P P - 1

2 0 0

180

160

140

1 2 0 (1) 0 c 100

(b :3 65 > I 01 I--

8 0

6 0

4 0

2 0

OI 0.oo

dTg = 55.4 x (fractional mol.wt. Increase) + 23.4 Rval. = 0 .9993

A

0 §

!

0 . 6 0 ...... I , I

1.20 1.80 _.1,

2 .40 3 .00

Frac t iona l mol.wt, increase

CHARACTERISATION OF POLYKETONE POLYMER SYSTEMS

BY THERMAL ANALYSIS TECHNIQUES

CHAPTER 9

297

CHAPTER 9: CHARA~RISATION OF POLYKETONE POLYMER SYSTEMS BY TRERMAL ANALYSIS TECHNIQUES

9.1 Introduction

A unique catalyst invention at the Shell Research Laboratories in Amsterdam in 1982 [I, 2, 3] made it possible to polymerise carbon monoxide and alpha-olefins such as ethylene into linear, perfectly alternating structures. This led directly to the development of a new class of thermoplastic polymers known as aliphatic polyketones (PK), which Shell is commercialising under the trademark CARILON.

Aliphatic polyketones based on carbon monoxide and ethylene are called PK copolymer, while the first commercialised grades based on carbon monoxide, ethylene and a small amount of propylene are called PK terpolymer.

It was exciting to follow the development of this polymer from the very first beginning on the laboratory bench up to a product ready to enter the engineering polymers market. This commercial significance also makes it a nice example to illustrate how different TA techniques (DSC, DMA, TMA and thermo electrometric analysis) focussed on one product contribute to the characterisation of such a new polymeric system. All data given are measured on non-stabilised, development phase PK co- and terpolymer samples made more than five years ago. Hence, these data can be different compared with those of the present, further developed, commercial grades.

9.2 Investigation of the crystalline phase of PK co- and terpolymers by DSC

9.2.1 PK CoDolvmer and PK terDolvmer The product-obtained after washing and drying of the reactor product is a white, semi-crystalline powder soluble in only very few exotic solvents i.e. hexafluoro-isopropanol (HFIPA) and meta-cresol. The presence of a crystalline phase in PK copolymer, as indicated by X-ray diffraction (XRD) analysis, is confirmed by a clear fusion effect measured during heating the sample in the DSC showing a Tm-value of about 258 C and a heat of fusion effect (Hfl-value) of about 152 J/g.

It is well-known that the occurrence of chain defects, in the form of for example small methylene sidegroups, could reduce this Tm-value [4]. This offered the possibility to reduce the relative high processing temperature of PK copolymers. The effect of addition of small amounts of propylene to the carbon monoxide/ ethylene mixture on the Tm-value is shown in Figure 9.1. A nearly linear decrease of the Tm-value as a function of the weight percentage of C3 was found for propylene concentrations between 0 and about 15 %wt. i.e. for the Tm- value holds:

298

Figure 9.1 Tm-value depression of PK copolymer due to C3 addition

5 5 0 . . . . . . . . . . . . .

54O -\Tm, K = - 5 . 9 x (%wt. C3) + 5 2 9

5 3 0 -

5 2 0 >

: s

~ - Tm(calc.) + 5K 50O

> I E I-- 4 9 0

Tm(ca lc . ) - 5 K .- !

4 7 0

4 6 0 ~ , ~

4 5 0 , ,~ 0 2 4 6 8 10 12 14

C3-con ten t , %wt.

299

PK terpolymer Tm(calc.), (K) - 529 - (5.9 x C3, %wt.)

The scatter in the with equation 9.1 calculated Tm values proved to be + 5 K.

9.1

9.2.2 The Tm(o)- and Hf (max) -values of PK CoDol_vmer Three extrapolation methods are, according to Cheng [5], available to determine the equilibrium melting temperature [Tm (o) -value] of a polymer:

- data for small molecules (low molecular mass homologs) are extrapolated to macromolecules,

- melting temperatures are extrapolated as a function of the crystallisation temperatures and,

- small crystal melting points are extrapolated to large crystal melting points.

The first method mentioned was used to determine the Tm(o)- value of PK copolymer. Extrapolation of the fusion temperatures of low + molecular mass homologs is possible using the relationship derived by Hay [6] between the Tm-value of oligomers and their degree of polymerisation (n):

Tin-value, K = (-2.R. [Tm(o)]z/Hf) . (Ln[n])/n + Tin(o) 9.2

The Tm-values of a series of oligomers are plotted as a function of (Ln[n])/n, extrapolation of (Ln[n])/n to zero results in a value for Tm(o). Five oligomer samples, with the chemical structure :

CH3 - CH2 - C0- (CH2 - CH2 - CO) x- CH2 - CH3

were synthesised with the x values 2, 3, 5, 6 and 8. The degree of polymerisation gives the number of basic monomeric units in a macromolecule i.e. for x = 2, for example , n = 7. The results of the fusion measurements on these oligomer samples are listed below-

Table 9.1 Results of DSC fusion measurements on PK copolymer, experimental oligomer samples

........... ,, J,,rl I re

" n -

~e value

13

15

7 J

9

19 i i,

f , ~r f i r

Tm- Hf- value, oC. value, j/g

83 185 l _ ,

116 192 ,,

156 200 , l, ,, l

170 217 , ,

190 222 q i "' , I,L • , ,'

300

The Tm-values of these samples increase, as expected with increasing values for x. These Tm-values are plotted as a function of (Ln[n])/n in Figure 9.2. Linear extrapolation to (Ln[n])/n is zero results in a Tm(o)-value of 599 K. Subsequently, using the slope value of - 868.8 K and the Tm(o)-value of 599 K, a Hf(max.)-value of 246 J/g was calculated.

The crystallinity, the x(c)-value, of a polymer is usually determined by X-Ray Diffraction (XRD) analysis. The ratio of the integrated crystalline intensities and the sum of the integrated amorphous and crystalline intensities is giving the x(c)-value. The XRD x(c)-values and the DSC Hf-values were measured for a series of PK copolymer samples. This permitted calculation of the Hf-value of completely crystalline PK copolymer, see Table 9.2 :

Table 9.2 Results of XRD/DSC crystallinity measurements on PK copolymer samples

l

x(c) -value (XRD)

0.65 , ,

0.66 ,,

0.58 _ _ , , ,

0.59 ' IUl

I I IIIIII

Hf-value, DSC

J/g

157

144

146

144 I I

III

Hf-value, x(c) - 1.0

J/g

242

218 ,,, ,,

252 . . . .

244 I

The by XRD determined average Hf-value for 100% crystalline PK copolymer is thus 239 J/g. This value reasonably agrees with the value of 246 J/g as determined by extrapolation of the oligomer results. Hence, the following basic properties for PK copolymer were obtained:

PK coDolvmer -- _

Tm (o) -value , K

Tin-value (DSC, 20~ K = 599 (326~ -- 531 (258~ _+ 1 (n = II)

Hf (max.) AHf ASf ASf (bond)

, J/g - 242 , kJ/mol - 13.6 , J/mol - 22.6

, J/mol - 7.5

Hf-value (DSC, 20~ J/g i.e. crystallinity x(c) ,

- 152 + 6 (n = II) = 0.63 _+ 0.03

Lommerts et al. [7] report a heat of fusion value between 215 and 330 J/g (12.0 kJ/mol and 18.5 kJ/mol) also based on extrapolation of Tm-values of PK oligomers. They are using,

301

Figure 9.2 Ln(n)/n - Tm relation for PK copolymer oligomer systems

6 2 0

0 d~

"(3

:::3

> I E

F--

0 03 a

5 9 0

5 6 0

5 3 0

5 0 0

4 7 0

4 4 0

4 1 0

38O

3 5 0 0.00 0 .05 O. 10 O. 1 5 0 .20 0 .25 0.3

Ln(n)/n

302

however, the Tin-value of a highly drawn fibre sample (i.e. 551 K) as Tm(o)-value. Allen et al. [8] report a heat of fusion value of 224 J/g (12.5 kJ/mol) for completely crystalline PK copolymer based on PK copolymer/ethylene glycol melting point depression measurements.

The crystallinity of a series of PK terpolymers was also determined by XRD and DSC to determine the Hf (max.)-value. The following basic values for a PK terpolymer system were obtained:

PK terDolvmer system Tm-value-(DSC, 20~ K - 493 (220~ _+ 3 (n = I0)

Hf (max.) , J/g = 207

Hf-value (DSC, 20~ J/g i.e. crystallinity x(c)

= 111 + 10 (n = 10) = 0.54 _+ 0.05

9.2.3 AiDha- and beta-crvstallinitv in PK CoDolvmer The PK copolymer chains can crystailise under certain circumstances into two different modifications i.e. the u-form and the E-form. Lommerts et al. [7] deduced from the observed X-ray reflections an orthorhombic unit cell for the u-form with a = 6.91 (2) ~, b = 5.12 (2) ~ and c = 7.60 (3) ~; the calculated crystalline density is 1.383 g/cm3. Chatani [9] reported for the E-form a, b and c values of respectively 7.97, 4.76 and 7.57 ~; the calculated crystalline density is 1.297 g/cm3.

All carbonyl dipoles in the u-form crystallites are pointing in about the same direction at equal height z. The dipoles of the carbonyl groups of the corner and the center chains are pointing in different directions for E-form crystallites. This difference makes the packing of the u-formvery effective. Lommerts et al. [7] report a cross-sectional area of the unit cell perpendicular to the fibre axis of 35.2 ~ while this value increases to 37.9 ~2 for the E-form material.

The u-phase crystallites are stable up to maximally 120~ The u-form crystallinity changes into E-form crystallinity at higher temperatures. This u/E crystal transition is indicated as the Tm'-value of PK copolymer. The fusion effect at about 258~ which is important for the processing of the polymer, is thus fusion of E-form material only.

Polyketone containing mainly u-form crystallinity at room temperature was reported by Lommerts et al. [7] for drawn fibre samples. Both Lommerts and Klop [7, 10] concluded that low molecular weight as-polymerised oligomers crystallise at least partly in the u modification whereas virgin high molecular weight PK copolymer (polymerised under the 'standard' conditions described in the patents [1, 3]) crystallises almost completely in the E modification with a

303

smaller crystalline perfection. Experiments with different polymerisation liquids showed, however, that under certain circumstances PK copolymers are obtained containing both u- and K-crystallinity. K-Crystal phase material also changes (partly) into u-crystal phase material due to high pressure and/or shear forces at 20~ Moreover, a proper annealing procedure proved to promote the presence of u-phase crystallinity in virgin high molecular weight PK copolymer

Table 9.3 Effect of annealing at T(isothermal) of a PK copolymer sample on the temperature location (Tm' -value) and the strength (Hct-value) of the u/E crystal transition

,,,, ,, ,,,,,

T- isothermal* oc

230 ,,

240

245 ....

250

255

260 . . . . .

265 ,,,, J,

Tin' - value, ~ . . . .

96 ,

108

115

120 , i

120 ,, ..... ,

96

91 ' I ,, ...... I

III I,,I!

Hct-value, J/g

2

5

8

12 i ,,,

15 , ,

4

1 , , I

* DSC procedure- heating from 20~ to T-isothermal, sample six minutes at T-isothermal, followed by cooling to 20~ heating from 20~ up to 275~ heating/cooling rates 20~

samples, see Table 9.3. The nummerical values in Table 9.3 and the transition effects in Figure 9.3 clearly show that both the temperature location and the strength of the u/E crystal transition go through a distinct maximum. One minute heating at 250~176 proved to be sufficient already to obtain a maximum amount of u-crystallinity in an originally only E- c~istallinity containing powder sample. Figure 9.4A shows the XRD spectrum of the non-thermally treated powder sample (only E-crystallinity). The XRD spectrum in Figure 9.4B is measured on the same sample after a thermal treatment of five minutes at 250~ Now, clear u-crystallinity reflections are present; the u/K ratio of this sample proved to be about 70/30. Thus, next to polymerisation liquid induced and pressure/shear forces induced u-crystallinity, thermally induced u- crystallinity is possible. This third method proved to offer PK copolymer samples with the highest u/K crystallinity ratio's.

304

o.,o ,SC'h..,.,,,g.:,. illi\~'r. 255c 0.55 2 6 C t m t n u t e . -

0.50

O. 45

~__ o.lo Ta 2 4 5 C /

1 ~ . / / \ \ = o.,, t ,,,..,.-" / \ ",,.

t __..- / \ - -

0 t0 ~ - _~,~.,,'- - \ ,

o.~-L . . . . . . . . . . . . . . . . ~;-':-;: ....... --.~., - z 3 . C . . . . . . . . . . . . . . . . . . . . . 0.00 I " . . . . " " " " ~ " ' * " ~ " - ~ " ~ " ~ 1 " 1 I I I " ~ ' ~ l " " " ' ~ " ~ " i ' ~ " = ' ~ ' l " " " ' " - " - " ~ " ' ~ ' ' * ' " ~ " ' ~ ' ~ I

50.0 70 0 go.o iiO 0 i30.0 i50 0

Temperature (~

0.50

0.45

0 40

A 0 3 5

0.30 3= o o ~ 14.

0 20

0 15

0 . t0

0.05

D S C h e a t i n g ==can -~/'" 2 8 C l m t n u t e / ! Ta " 255C

/ ~ / I

/ i / . / i

/ , l \

~ i

. _ . . . t " "X.

. . . . . . . . . . . . . . . . . . . . a = 2 6 O C " ' ~ " " . . . . .

. 1 I , ~ , T a = 2 6 5 C

, . ~ . . . . ,..,,,. - - - . ~ . ~ - ~ * ~ - " ~* '~ ~ . ~ * ~ . . ~ . . , , . .~ ...... # . , . . . . .~ . ~ . p . ,~,. ~ . ~ , . . ~ _ ..e~ - . ,~

0 0 0 - I I I i I I I I I 50 0 70 0 g0.0 i t0 0 t30 0 i50 0

Temperature ( ~

Figure 9.3 The alpha/beta crystal transition of a virgin PK copolymer sample after annealing (6 minutes) at the indicated temperatures Ta

1 .20 305

1.08

0 .96

0 . 8 4

0 .72

O. 60

O. 48

0 .36

0 .24

0 .12

1 0 . 0 20 .0 30 0 40 0

I .00

Figure 9.4B Thermally treated sample alpha and beta (70/30) crystallinity

II i i , , i

50 .0 60 .0

0 .90

0 .80

0 .70

0 .60

0.50

0.40

0.30

0.20

O.iO

10.0 20 .0

I I Figure 9.4A only beta crystallinity

ii I non-thermally I I I I treated sample

30.0 40.0 50.0 60.0

Figure 9.4: XRD spectra of PK copolymer samples

306

Polymerisation liquid induced and pressure/shear forces induced u-crystallinity effects proved to be irreversible (DSC) or partly irreversible (XRD) after heating through their u/g-crystal transition. Additional DSC experiments were performed to investigate this aspect for the thermally induced u-crystallinity.

A sample with a thermally induced u-crystallinity phase was heated in the DSC to respectively 140~ 250"C, 270~ and 280~ The relative strong u/g-crystal transition proved to be completely reversible (see Figure 9.5) after heating to temperatures between 140~ and 250"C. First after heating to 270"C, the u/E-crystal transition is strongly reduced and is shifted to lower temperatures. The results in Figure 9.5 also show the g/u crystal transition at about 50~ during the cooling scans from 140~ and from 250~ down to 20~ The recrystallisation effects become small i.e. non-detectable by DSC, after heating up to temperatures �9 250~

The optimum u-crystallinlty promoting annealing temperatures proved to be 250"C/255~ (Table 9.3) i.e. in the melting region of the g-phase. Thus, the E-phase crystals might also be influenced by this annealing procedure. The results of fusion experiments showed that both the Tm-value and the Hf- value of the g-crystal phase increase due to the annealing procedure. The following average values were measured for the same PK copolymer powder sample:

Tm-value, oc

non-treated powder samples: 258 • 1 (DSC first heating scan)

therm, pretreated samples: 264 _+ 3 (u/g transition at 120~

Hf-value, J/g

152 + 6 (n = 11)

166 + 4 (n = 4)

These results indicate that the perfection or the size of a part of the g-crystallites is improved/increased during the annealing procedure at 250~176 Only the g-crystallites with a certain degree of perfection or a certain minimum size are able, subsequently, to transform from a g-structure into the u-structure during cooling to temperatures below 50~

307

I o

e l

I 1

I . i . ~,, �9 �9

~ l i p ,

~--- e e .

I i t .

t l .

I . I

I . t .

IO ,

~ ) I I .

t .4,

I I ,

0 ,.,.

' *~ e l . (D 3: "i

e

t 4 0 C

� 9 I t n o l o i n I �9

. . . . . . . �9 .J . ' , ,~ , ' ,~ , "" ,.~o ' ~ . '

Temperature (~

% ( - - . , I " ' ~ ' " "

~; COO 1 I n9" ,~ ,

s ~ l l n 2 . , , . . . . . . �9 ,4 . ~ o " ,,,;o ' .,,~.

Temperature (~

) l i t

I O t . . . . . . . . . . . . . . . . . . ~ ' , . o., / }

I I �9 . . "

"F,,,...._ - - "

"~ " l ~ - " " " ~ =oo w,,.,o " 1 " 0 . 4

::J,~ '- , i = . ~ o ,do ' . , , ' ~ . = . ,

Temperature (~ I . I ] . . . . . . . . . . . I . . . . . ~' -

1 .1" /

I . l t s .~ / / h i t l ~ I n I . - - "

: " ' L 2 0 0 C

"'1 - - - - " I . I

l * �9 s c l n c o o I I n o O l " j ' I i ! - . ~ " i " ' 1 i i �9 - - i ~

I I I t O M i l l I I I I P i l l ~l~O

Figure 9.5 T e m p e r a t u r e ( ~

Results of subsequent D S C heating/cooling scans on a virgin PK copolymer sample

I I 2 5 8 C

I I I I !

", i ; % J ; . . . . . . . . k

I t o e N e e

308

9.2.4 Alpha- and beta-cry_s~llinity in PK GQpol_vmer ~fter a common processinu procedure

During the most frequently used processing procedures i.e. injection/compression moulding, the polymer is first completely fused. That implies for PK copolymer that the K- crystalline phase first completely is destroyed, before development of any u-crystallinity due to certain annealing conditions can be expected.

This fusion/recrystallisation process determines the degree of perfection of the E-crystalline phase formed and this (as we saw above) partly determines the efficiency of an annealing procedure ment to promote u-crystallinity in such an article. The extent of such an effect was first investigated by a series of DSC experiments, see Table 9.4.

The experimental results in Table 9.4 indeed show that the optimum annealing temperature shifts from 250~ for a non-fused powder sample to about 225~ for a first sample. Annealing during less than one minute at 250"C/255oc was already sufficient for the formation of an u-crystalline phase with a maximum Tm'- and Hct- value in a non-fused powder sample. This time increased for the fused samples from less

_

than one minute into about six minutes. Moreover, the maximum Tm'-value reached now proved to be 106~ instead of the 120~ measured for the powder sample. This illustrated that after recrystallisation from the melt either less perfect E- crystallites or smaller E-crystallites are obtained than after crystallisation from a liquid. The annealing process is clearly not able to eliminate this difference completely.

Table 9.4 Effect of annealing at T(isothermal) of PK copolymer on the temperature location (Tm'-value) and strength (Hct-value) of the u/E-crystal transition (totally fused sample series)

, 1 1

T(isoth. ) * oc

180 ,,

200 , ,

210 ....

215 ;

220

225

230

240

250 I

rl

Tin' -value, oc

97

98

99

I00 � 9 ,

103

106

94 ,

93 ,

92

,i

Hct-value, J ' /g

10 , ,

8

10

8

8

Ii

4

309

Table 9.4 continued.

* DSC procedure- heating from 20~ to 275~ cooling to T(isothermal), sample six minutes at T(isothermal) followed by cooling to 20oc, heating from 20~ to 275~ heating/cooling rates 20~

Subsequently, the effect of fusion before annealing was investigated using a series of compression moulded samples. First a one millimeter thick sample sheet was prepared by compression moulding during 2.5 minutes at 280~ followed by cooling as fast as possible (water-cooled mould) to 20~ Two other sample sheets were prepared in the same way, but first placed in a second mould heated at 240~ during respectively 6 minutes and i0 minutes before cooling to 20~ too. These samples were analysed subsequently, see Table 9.5.

Table 9.5 Introduction of u-crystallinity in compression moulded PK copolymer sheet material by an annealing procedure

- r ,

measured property

. , , ,,

u-phase Tm' -value, "C Hct-value, J/g u/E ratio, estimated

E-phase i Tin-value, oC Hf-value, J/g x(c) density, g/cm3

die1. constant tan delta

, , , , �9 JL I

" , llrll III ' ~,.,

PK co- compr. polymer moulded virgin 2.5 rain. sample , 280" C

n ~

n.p. o/~oo

2.5 min. 2.5 min. 2 8 0 ~ + 280~ +

6 min. 10 min. 240oc 240oc

106 112 5.1 8.3

25175 40160

253 251 119 116 n.d. n.d.

,.

1.284 1.299

n o p .

n.p. o/loo

258 152 0.63

n.d.

255 111 0.46

1.267

n.d. 5.30 n.d. 0.0171

n.p. -- not present, n.d. = not determined, properties measured at 22~ dielectric properties measured at 1 kHz.

5.08 4.99 0.0150 0.0127

[II Ir

Lo~nerts [7] reports for u-phase PK copolymer a density of 1.383 g/cm3 and for E-phase PK copolymer a density of 1.297 g/cm3. Besides, he assumed a density value of 1.21 g/cm3 for

310

the PK copolymer amorphous phase material. The compression moulded reference sample (i.e. 2.5 min./280~ consisted next to an amorphous phase of an fully E-modification crystalline phase. The crystallinity of such a polymeric system can be calculated according to-

x(c) = [p - p(a ) ] / [ p (c ) - p(a)] 9.3

where: x(c) = crystalline fraction i.e. 0.46, p = sample density i.e. 1. 267, p(c) - density of E-crystalline phase i.e. 1.297, p(a) = density of amorphous PK copolymer,

Using the mentioned values, an amorphous phase density of 1.241 g/cm3 was calculated, which seems to be an acceptable value. (The x(c) values of the virgin powder sample and the sample after compression moulding at 280~ (2.5 minutes) can be calculated using the Hf(max.)-value of 242 J/g (see 9.2.2) because both the DSC and the XRDmeasurement were performed on PK copolymer with only an amorphous and a E-crystalline phase. This XRD value does not hold, however, for PK copolymer with both an u-crystalline and a E-crystalline phase).

The measured density values increase due to the annealing procedure and this increase agrees with the increase of the estimated u/E-ratio.

Both the dielectric constant and the dielectric tan delta values are (partly) depending on the possibilities to move on molecular scale, see 5.1.2. Hence, the level of both dielectric values has to decrease with an increase of the density and this effect is confirmed by the experimental values in Table 9.5.

Six minutes isothermal at 240~ is a condition which will be rarely met during common processing operations of PK copolymer systems. Hence, the u/E-ratio of PK copolymer articles after 'standard' processing operations will be low. The u/E-ratio strongly increases, however, if high shear rates are present during the processing procedure. Klop et al. [10] reported that high molecular weight PK copolymer with an almost completely E crystalline phase changed due to a process of melt-spinning into nearly completely u crystalline material.

9.2.5 Alpha- and beta-crystallinitv in PK terpol_vmers Klop et al. [10] showed that for PK terpolymer fibre samples, the b and c dimensions of the E-phase unit-cell insensitive are to changes in the C3 concentration. The a dimension, however, clearly increases with increasing C3 concentrations. This suggests, according to Klop, that the methyl groups are incorporated into the PK terpolymer E-phase lattice as defects. They also calculate for an uniformly random distribution of methyl groups along the polymer backbone chain, an average distance between in-chain neighbouring

311

methyl groups of about 54 ~ for a PK terpolmer with about 7 %wt. C3. The crystallite size in the fibre axis direction ranged from 120 to 200 ~. Hence, the methyl groups can indeed be incorporated into the E-phase crystal lattice. It will be not surprising therefore, that the presence of methyl groups hampers the formation of u-phase crystallinity. The results shown in Table 9.6 confirm.this idea. These data show that for PK terpolymer systems with a C3 content ~ about 5 %wt. hardly no u-phase crystallinity can be induced by a thermal treatment.

Table 9.6 Thermally induced u-phase crystallinity in PK terpolymers.

, , T ' , L

PK terpolymer C3 cont., %wt.

0.0

1.24 , , .

1.5

3.7 I I I "

, r r I , , , , , , f ' , , , I Ir ,= I

annealing, 5 rain. at T isoth., ~

. . . . . .

255 ...................

248

245 ........

243 . 'I I It' r ........

T ~ l -

value "C

,

123

106

98

85 ..... ,.,,

Hot- value,

J/g

2O

Ii

9

3

Summarising, PK copolymer chains can crystallise in an u- modification and in a E-modification. The crystalline phase of virgin low molecular weight oligomer systems is at 20~ mainly in the u-modification. The crystalline phase of high molecular weight virgin PK coploymer, on the other hand, is (again at 20~ completely in the E-modification. PK copolymer systems containing both crystal modifications are obtained by-

- the use of different liquids/solvents during the polymerisation process,

- cold compression of virgin reactor samples, - a proper annealing procedure, - compression/shear forces in combination with the thermal

treatment during processing operations like compression moulding and gel-spinning.

Besides, u-phase crystallinity was found to be present only in PK terpolymer systems with C3 contents less than about 5 %Wt.

For PK copolymer u-phase crystallites, all carbonyl dipoles are pointing in about the same direction at equal height z, whereas in the E-phase crystallites the dipole of the cabonyl group of the corner and the center chain point in different directions [7]. This might mean that the presence of a strong electric field during the recrystallisation from the melt of PK copolymer also might promote the formation of u-phase crystallinity. A high u/E-crystal ratio might improve the strength/stiffness properties of PK copolymer but perhaps the most important effect might be an improvement of the barrier properties.

312

9.3 Investigation of the amorphous phase of PK terpolymer by D~/DSC

9.3,1 Amorphous Dhase transition effects Dynamic mechanical analysis (DMA) on both a cold (20"C) compressed powder sample and a compression moulded (240~ sheet sample was used to investigate the relaxation processes i.e. especially the glass-rubber transition, of the amorphous phase of PK terpolymer.

Figure 9.6 shows the (shifted along the E" axis) dynamic loss modulus (E")/temperature curves of the cold compressed powder sample and that of the compression moulded sheet sample. The crystallinity of the cold compressed sample is about 57%, while the crystallinity of the sample after compression moulding at 240oC i.e. after recrystallisation from the melt, is decreased to about 42%. Both samples are measured at a frequency of 1 Hertz and a heating rate of 2"C/minute. The relaxation behaviour proved to be clearly changed due to this difference in crystallinity of about 15 %.

The higher crystalline, cold compressed sample shows a so- called crystalline phase (u) transition at about 130~ a (weak) glass-rubber (E) transition at about 50~ and a secondary, amorphous phase (Y) transition at -75~ This weak glass-rubber transition effect is typical for a semi- crystalline polymer. It indicated already that it would be difficult to detect this effect by DSC.

The (crystalline) u-transition has been disappeared completely for the lower crystalline (at 240~ compression moulded) sample. The intensity of the glass-rubber (g) transition, however, has been clearly increased. Moreover, the g- transition E" (max.) temperature is shifted from about 50~ to 15~ The Tan ~ (max.) temperature is used to indicate the (DMA) Tg-value of compression/injection moulded PK co- and terpolymers. This Tg-value proved to be 19"C _+ 2~ with no significant difference between these values measured for PK copolymer and PK terpolymer.

The y-transition in Figure 9.6 is the only transition effect which seems to be not influenced by the difference of about 15 % in crystallinity between both DMA samples.

9.3.2 Ageing and moisture absorption effects Two effects were noticed during the determination of the mechanical properties of PK terpolymers: - the flexural modulus increased as a function of time during

storage at 23~ and 50 % R.H. and, - a weight increase of the test samples due to moisture

uptake. Soon became clear, however, that a short thermal treatement (15 minutes/140~ was sufficient to restore the original modulus value.

7 . 0

A m

a .

. 1

o ) 0,

7 . 3

8 . 8

8 . 5

Figure 9.6 The dynamic loss (E") / temperature curves of a PK terpolymer sample after room temperature compression and after compression moulding at 240~

' ' ~ s : �9 S �9 j .. : - .

. . . .-" V

POWDER SAMPLE COLD PRESSED AT 28 deg. C . o ~ , ' * " ~ �9

,,.4 % �9 �9 ~r �9

e

�9 �9 e 0 �9 �9 :cO e

" t " t e e

�9 ~.~,.. �9 SAMPLE COMPRESS I ON ~,,,,r " MOULDED AT 24~ ~

% " t

� 9

""2..%..'" .. -o.~,,~,,',, t . " e e e ' l i ~ �9 o o - "~ �9 �9 �9

6

I . I, . I I , , I I ~I . I I ! !

~ ~ I ! ! t

%

8.25

A w

a .

' e OD 0

_1

.. ?.75

�9 7 . 5

. t I . l i ! f . I . l I l I I I i I * i i i r i i 7 ' , 2 5

= m = = | | m = J | ~ Temperature (~

314

A modulus value increase upon storage under ambient conditions is also reported for other semi-crystalline polymers like, for instance, polypropylene. Struik [11] measured for PP a continuously increasing dynamic stiffness at 20~ in combination with a decrease of the intensity of the glass- rubber (8) transition of PP (the temperature location of the g-transition did not change). Struik called this phenomenon an amorphous phase ageing effect; a densification process of the amorphous PP phase due to a free volume relaxation effect.

This explanation was thought to hold also for the modulus increases measured for the polyketone systems. Two 'aged' samples were, therefore, investigated by DMA:

- one sample stored during 720 hours at 21~ in an exsiccator i.e. a moisture free sample and,

- one sample stored during >1500 hours at 21"C and 50 % R.H.

Both samples were measured in a Polymers Laboratory DMA with a frequency of 1 Herz and a heating rate of 2"C/minute from - 100"C up to 150"C in a nitrogen atmosphere. Possible ageing and absorbed moisture effects are removed during heating to 150~ The first DMA scan was, therefore, immediately followed by a second DMA scan to measure the same properties but now on the non-aged, dried sample. The DMA results in Figure 9.7 of the dry, aged sample show an increased dynamic stiffness at ambient temperatures (Figure 9.7B) and a decrease of the intensity of the glass-rubber transition (Figure 9.7A). These results confirmed the idea about amorphous phase ageing of polyketone polymer systems especially because no secondary crystallisation effects could be detected.

The sample stored at a relative humidity of about 50 % contained an equilibrium moisture saturation of 0.74 %wt. The Tan 6 maximum of the ~ relaxation in Figure 9.8A is now not only decreased due to the ageing effect but also shifted from 19~ to 5"C due to the plasticising action of the absorbed moisture. The dynamic stiffness at 20~ of the aged sample (Figure 9.8B) is now IEW~ than that of the non-aged, dried sample due to the plasticising action of the absorbed moisture. The same sample was, subsequently, stored during seven days in ion-free water to reach an equilibrium water saturation of 2.33 %wt. The Tg-value shifted, due to this absorbed water phase, from 19~ to -8~

It will be clear that the mechanical properties of polyketones at ambient temperatures are sensitive for these ageing and these moisture absorption effects especially due to the presence of the glass-rubber transition in that temperature region. The influence of both effects is in general opposite to each other; the stiffness increases due to ageing and decreases due to moisture absorption. Moisture absorption effects are time and object dimensions dependent whereas ageing effects are only time dependent.

315

@ "o s

o)

Q rn

aged =amp le

n o n - a g e d 8 amp I e ( a f t e r ' h e a t t n g t o 1 5 0 ~

" i l l .

4 0 -tO O tO m N 4 U m ~ N U Temperature (~

FIGURE 9.7A

|N

0.41-

0.!1. m I1.

O.Ib

UJ |.11" r

_o .c "o.o. c

an.o,

| . 0 .

e d =tamp l e

0.7. r

-Ill '-tO O tO m m a U U M N Temperature (~

FIGU~d~ 9.7B

The results of the DMA m e a s u r e m e n t s on a PK terpolymer sample after 720 hours of ageing in an exsiccator under vacuum and a temperature of about 21 ~

316

. g ,

~ . .(r/�9

" 0 C:

O) r : 6 .0 l . r - Q

irTt __

aged eemp le

Sample a f t e r drytnS~ (by hea t ing to LSBuC) (ue tght loss B.?4 ~.)

- m ~ - IO 0 tO N N ~ B N 70 N m

Temperature (~ FIGURE 9.8A

I N

11.4.

I I . J

0 . 1

UJ r e . t

o 0') C O.O, ..~

" 0

O.O-

=amp I

O.Y.

; , , , , , , , ; , , , ; ,;, ~ -W - - I0 @ tO I0 30 40 N IO

FIGURE 9.8B Temperature (~

The results of the DMA measurements on a PK terpolymer sample After > 1500 hours ageing under room temperature conditions

317

Figure 9.9, finally, shows the Tan 6/T curve of the aged DMA sample with an equilibrium moisture saturation of 0,74 %wt. and the same curve after drying due to heating to 150~ The 7-relaxation, clearly present in the 'wet' sample, is strongly reduced in intensity in the dried sample. Hence, absorped moisture is for an important part responsible for the intensity of the y-relaxation of polyketone (such an effect of moisture absorption was described before, see 5.2.2). This might also be the reason for the nearly identical 7-transition effects as shown in Figure 9.6.

.on.

. W

.OT.

-o c

. N

==

. 0 4 .

.O~l,

.Or, t ' I . . . . . . v ~ v 1 u t I~

- I |@ ~ - 7 0 -SO - 3 0 - I 0 ! 0 30 SO 70

Temperature (~

S e c o n d l r y G l a s s - r u b b e r t r a n s i t i o n s t r a n 8 ! t t o n 8 ( b e t a - r e I a x a % t o n s ) ( g a m m a - r e | t l x a t ! o n s )

m o i s t u r e s a tu r-lLtecl / /~ \

_

FIGURE 9.9

The effect of a sample drying procedure (by heating to 150 C) on the low-temperature gamma relaxation of a PK terpolymer

318

9.3.3 Determination of the Tu-value of PK terpol_vme~ by DSC The first attempts to determine the PK co- and terpolymer Tg- values by DSC seemed to fail until we realised that the above mentioned ageing and moisture absorption processes might hamper this measurement. The Tg-value determination of PK co- and terpolymer by DSC was tried, therefore, using non-aged and dried samples. These DSC experiments were performed as specific heat (Cp) determinations on samples of about fifteen milligramme, using a heating/cooling rate of 20~ and the following temperature programme:

- heating from 20~ to 150~ 15 minutes at 150~ to dry the sample and to remove ageing effects,

- heating from -100~ to 250~ to measure the Tg-value and the first fusion effect i.e. after crystallisation from solution (scan i),

- heating from -100~ to 250~ to measure the Tg-value and the fusion effect after recrystallisation from the melt (scan 2).

The Cp/temperature curve of the non-aged, dried powder sample is shown in Figure 9.10A. Now, a weak glass-rubber transition effect (onset 6~ is indeed visible in the Cp/T curve.

The crystallinity of the, from solution crystallised, powder sample is 57 %wt. (based on a Hf-value of 207 J/g for I00 % crystallinity, see 9.2.2). The crystallinity decreases, due to recrystallisation from the melt during the second scan, to 43 %wt. (Hf-value 88 J/g). Figure 9.10B shows that the resulting increase of the amorphous phase is reflected in a stronger glass-rubber transition effect (onset value 3~ Both Cp/T curves are plotted on an enlarged scale in Figure 9.11. The glass-rubber transition effect of the second Cp/T curve, measured after recrystallisation from the melt, approximates the for a glass-rubber transition characteristic step-wise Cp/T change. The glass-rubber transition effect of the first scan, measured after crystallisation from solution, is that of an amorphous phase strongly influenced by the presence of the crystalline phase. Such an effect was already detected with the DMA experiments, see 9.3.1.

PK co- and terpolymer Tg effects can be detected by DSC both on dry, non-aged virgin powder and on fused samples. The results of the DSC Tg(onset)-value determinations on (virgin) powder samples proved to scatter, however, considerably due to the too strong influence of the crystalline phase on the shape of the Cp/T curve. Thus, a real DSC Tg (onset) -value determination of PK co- and terpolymer samples is only possible on non-aged, dry and one time fused samples. The reproducibility of this DSC Tg(onset)-value determination proved to be _+ 3~ The average DSC Tg(onset)-value of a series of PK co- and terpolymer samples proved to be 4"C.

The ~ crystallinity of the sample in Figure 9.11 after recrystallisation from the melt is also reflected in the hiuher Cp values in the temperature region between Tg and Tm,

g.O

8 .0

3.O

0 o 7 . 0

6* 6 . 0

0 -r. 5 . 0

E ~ 4 .0

2.0

g .o

I . O - [

- I 0 0 . 0

l ' Tm-value = 2 ' 2 3 ' ~ .Hf-value = 119 J/g j

=

Tg-value = 6 deg. C I (DSC onset) [ ~ 2 / ..... ,

....

319

I I I I I ' - I " " 1 -50 0 0 0 50 0 tO0 0 t50 0 ;~00 0 2 5 0 . 0

RATE1: 20~ Temperature (~

FIGURE 9.10A The (DSC) Tg effect for a PK terpolymer dried, non-aged powder sample

Tm-value = 217 ~ Hf-value = 88 J/g i.e. x[c] = .

B.O

O) ~,, 7 0 - 3 v

6 . 0 -1- ._o ~.o 0

~ - 4 .0

3 . 0

2 0

t O /.- ..

Tg-value = 3 deg. C (DSC onset)

�9 ' I " 1 ...... I I . . . . I " - - - -+1 - t O 0 . O -50 0 0 0 50 0 100 0 150.0 200 0 ZSO.O

R A T E 1 " 2 0 ~ Temperature (~

FIGURE 9.10B The (DSC) Tg effect for a PK terpolymer sample after recrystallisation from the melt

I ~ 2.5 "~

. X:

2.3 "-~ q) o. 09

2.1

t.9

1.7

t.5

t.3

t.1

.9

Figure 9.11 .... , , . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

The specific heat (Cp) of PK terpolymer as a function of the temperature (dried, non-aged sample)

+ non-fused, powder sample, Hf = 119 J/g

x sample after recrystallisation from the melt, Hf = 88 J/g

+--X -'+ --X - + . - •

X .// • /x.,. ; . .

Both dotted 'base, lines coincide, after extrapolation into the polymers' me1 tphase. The Cp-values at 240 C measured during both scans are 2.340 J/g.C

- ............

/ § / • +

/• +/ • / +

• /

+,

•

0

.7 cg

I

6~ C~ E~

I I I

ED E~ C~ CD oJ u3 r --.

9 - - 4

Temperature, ~

C~ ED ~

-. -. ~u 0u

see Table 9.7-

321

Table 9.7 Specific heat of PK terpolymer at temperatures between Tg and Tm.

temper ature,

oC

60

7O

8O

90

I00

110

I II ,,

Cp-value, J/g.C x(c) =0.57 scan 1

i

1.334

1.386

1.440 ,,,

1.495 ,.,

1.555

1.615 ,, If

Cp-value, J/g.C x(c)-0.43 scan 2

..

1.447 , i , ,

1.501

1.557

1.612 , ,,

1.669

1.730 I ' Ill

ACp, J/g.C

0.113 ,

o .,~5

0.117

0.117 ,, , ,,

i,w l

0.114

0.115 I

The average ACp value of 0.115 J/g.C is assumed to be the increase of the glass-rubber transition ACp-step due to a crystailinity decrease from 57 %wt. to 43 %wt.

The ACp-step at Tg measured on the sample after recrystallisation from the melt (scan 2) ~ is 0.323 J/g.C. The ACp-step at Tg during scan 1 can, due to the shape of the Cp/T curve, not be determined directly. It can be estimated, however, by subtracting the ACp increase between Tg and Tm from the ACp-step at Tg measured during scan 2 i.e. 0.323 - 0.115 - 0.208 J/g.C.

A third phase, next to the amorphous phase x(a) and the crystalline phase x(c), is thought to be present in the temperature region between Tg and Tm of semi-crystalline polymers which is called the rigid amorphous phase x(r,a). This is amorphous material hindered to such a degree by the crystalline fraction that it behaves rigid [12, 13]. Assuming that a ACp-step at Tg of 0.115 J/g.C corresponds with a polymer weight fraction of 0.14, the extent of the possible rigid amorphous phase in PK terpolymer can be estimated. The sum of x(c), x(a) and x(r,a) is 1.00, thus might hold for PK terpolymer (after recrystallisation from the melt):

0.43 + ([0.323/0.115] x 0.14) + x(r,a) = 1.00 9.4

x(r,a) = 0.18

Thus, PK terpolymer after recrystallisation from the melt might have a rigid amorphous phase (in the temperature region between Tg and Tin) of about 18 %wt.

322

9.4 TMA measurements on PK terpolymer systems

9.4.1 The linear thermal expansion coefficient of lona - - v

~lassfibre reinforced PK terDolvmer systems Glassfibre reinforcement is often used to modify the physical properties of thermoharding and thermoplastic polymer systems. The linear thermal expansion coefficient (1.e.c.) of a polymer is one of these properties which can be influenced (i.e. decreased) by glassfibre addition. In chapter 3.1.2 is shown that the 1.e.c. of a polymeric system at equal filler/ glassfibre concentrations decreases with an increasing average fibre length. The processing techniques for thermoharding polymers allow application of long glassfibres resulting in systems with low 1.e.c. values.

Long glassfibre reinforced (LGFR) thermoplastic sheets can also be prepared using the so-called wet deposition technique. First, the polymer particles and the about 10 millimeter long glassfibres are thoroughly mixed by preparation of a water based slurry. Subsequently, the water phase is removed by cold compression, followed by a compression moulding step at a temperature �9 Tm. The 1.e.c. reducing possibilities of this method for PK terpolymer was investigated with a series of samples with respectively 0, 5, 13.6 and 20 %v. of glassfibres.

Sample sheets of 100 x 50 x 4 millimeter were prepared with the wet deposition technique. Two rectangular TMA samples and one circular TMA sample were machined from these sheets. The rectangular samples, i0 x 7 x 4 mm, were taken in the length (X) direction and in the width (Y) direction of the sample sheet. The circular sample, diameter 5 mm, was used to measure the l.e.c, in the Z direction. The samples, placed in a TMA 7, were measured from -20oC to 120~ with a heating rate of 2"C/minute. Subsequently, the sample was cooled with the same rate and heated again for the real measuring scan. The first heating scan was ment to remove the frozen-in stresses which influence the thermal expansiDn behaviour (see 3.1.2).

Figure 9.12A shows that this stress release effect for the X- direction of the reference sample (no glassfibres) results in a permanent length decrease, just like the effect shown in Figure 3.1. The same effect was also measured in the Y- direction. In the Z-direction (Figure 9.12B), however, an expansion effect is measured. The extent of the measured shrinkage/expansion effects are listed in Table 9.8. These values show that these shrinkage/expansion effects are clearly caused by the compression moulding process; the measured effects decrease with increasing glassfibre contents.

The l.e.c, values measured during the second heating scan proved to be nearly equal for the X- and Y-direction. Figure 9.13A shows the l.e.c, as a function of the temperature for the

323

lO0. O

100.6

100.4

100.2

o~ I00 .0

-1- ~.~ g9. o C 0

9g. 6

Ic gg. 4 Q 0,,

99, 2

gg. 0

go. 9

J f l ee t

c h r I nkege e f f e c t

/ 7

s e c o n d coo l lns l asCam

. J /

. / /

.7 . /

. / /

, I , I I I I " 1 I

-~o o.o ~.o so.o ?s.o Ioo.o

T e m p e r a t u r e (~

Figure 9.12A The length change of a PK terpolymer compression moulded sample during heating/cooling scans (X-direction)

100.8

100,6

100.4

100.2

I00 .0

~ l .O

gg, o

~ 1 . 4

~ . 2

gg.o

gB.o

103.0

102.5

102.0 i -r-

c 0 . . = .

. s , . . .

101.5

c 0

I01.0

100. 5

I00. 0

. . . . . . . . . .

. . ~ ~

aeocnd cool Ing s c l n / ~ 1 7 6

expmns t on effect,

i l . . . . . I l ~ i I I I . . . . . zs. o o. o zs. 0 so. o ?5. o 1o~ o T e m p e r a t u r e (~

Figure 9.128 The length change of a PK terpolymer compression moulded sample during heating/cooling scans (Z-direction)

-lO'J. 0

-102, 5

-I02. 0

-tOt.

"101.0

100.5

100.0

15 .8

13 .5

12 .9

1 0 . 5

9 . 9

7 . 5

)

6 . 0

4 . 5

3..'

1.5

0.( ] Q

/

LIN . EXPRNSION COEFF. ( x ( - 5 , I / K )

+ r e f e r e n c e

s a m p l e

/ + ( n o ; ] 1 r o s e )

/ + / f §

/ +

/ + "

J 4-

Figure 9.13A Linear expansion coefficient of LGFR PK terpolymer composites as a function of temperature and glass-fibre content (Y-direction)

-- . - -- X ~

~ ' - x 5 Y,v. ; ] lass " ~ X ~ - - X - . . . - - X - - - - - X .- . . -- X - - . - - X

\ a \ 13 .6~v ;] I ass

~ o . . ~ . o _ _ . . . o . . _ _ . o . . . . _ _ o ~ O _ _ . . o 20 ~,v. g lass

m . l , I . l . I �9

T E H P d R R T U R E , d e ; ] . C

I �9 i . I . I _ . t . m . ! Q w m ~ (= m m (,D r ~ G3 01 m -- . N

,.4 ~ =.4

2;'

24

21

10

t5

12'

L IN . s COs ( x E - 5 , I / K )

4 / /x ql x

/x/ <1 /x/

+-/ + /

§

/

~ J /

X

/ 4 [3.6~.v ; ] l u =

4 x 5 %v ;]lass / X

+ / . + r e f e r e n c e

s a m p l e ( n o g I a s s )

Figure 9.13B Linear expansion coefficient of LGFR PK terpolymer composites as a function s of temperature and glass-fibre content (Z-direction)

3

T I E H P E R f l T U R E t d e ; ] , C

Q . n . m . m . , . i �9 I - i . , - m l _~.. I , . , L t = m m ~ = = m m ~ & m =

(.,J

r

325

Table 9.8 Thermal shrinkage/expansion effects of LGFR PK terpolymer systems due to heating to 120"C

fibre content

% v.

0

5

13.5

20

,, ' I I

X-direc -tion,

shrinkage %

1.08

0.42 , ,

0.19

0.20

I

Y-direc -tion,

shrinkage % i

0.99

0.44 ,

0.18

0.10

r i l l _J" III

Z-direc -tion,

expansion %

1.19

0.48 J ,,, ,

O.O7

0.20 , , ,'

Y-direction. The l.e.c, values decrease not only due to the glassfibre addition; the temperature dependency of the l.e.c. also decreases. The efficiency of the glass addition strongly decreases for volume percentages ~ 13.5 %. Figure 9.13B shows the same results but measured in the Z-direction. Both the level and the temperature sensitivity of the l.e.c, values increase in this direction due to the glassfibre addition.

These data illustrate that addition of long glassfibres indeed strongly reduce the l.e.c, values of PK terpolymer in the X- and Y-directions but at the cost of a considerable amount of anisotropy in the Z-direction.

9.4.2 The repeatability of the l.e.c, determination A large injection moulded PK terpolymer sheet sample (350 x 110 x 2.5 ram) was used to determine the repeatability of the 1.e.c. determination. Four 40 x 15 x 2.5 mm samples were taken out of such a large sheet, see Figure 9.14:

sprue

110 ram.

* . . �9 . �9

"15 ~ * * ,

350 mm~.

U-7

* * X-direction * * ~ Y-direction �9 * ~* ~- 15 ~m. * ,

, . , , . , �9 . , .

~ measuring direction

15 mm.

~ 9 40 mm.

Figure 9.14 The location of the TMA samples in the c.m. sheet

326

Table 9.9 Results of l.e.c, repeatability measurements on injection moulded PK terpolymer samples.

sample number

X-direction sample A sample C sample D

average values

"~ -- _ , ,,

Y-direction sample A sample B sample C sample D

average values

Z-direction sample A sample B sample C sample D

linear expansion coefficient l/K, 0~ 20~ 40~ 60"C

, ,=

average values I i

at 80oC

. . . . . . . |

8.85E-5 1.13E-4 1.21E-4 1.28E-4 1.36E-4 8.85E-5 1.09E-4 1.16E-4 1.23E-4 1.30E-4 9.31E-5 1.11E-4 1.17E-4 1.24E-4 1.31E-4

9.00E-5 1.11E-4 1.18E-4 1.25E-4 1.32E-4 • 0.27 • 0.02 • 0.03 • 0.03 + 0.03

..... , , , =

9.41E-5 1.06E-4 1.12E-4 1.18E-4 1.26E-4 9.00E-5 1.09E-4 1.17E-4 1.24E-4 1.32E-4 9.17E-5 1.11E-4 1.18E-4 1.25E-4 1.31E-4 9.02E-5 1.09E-4 1.16E-4 1.23E-4 1.31E-4

9.15E-5 1.09E-4 1.16E-4 1.23E-4 1.30E-4 • 0.i0 • 0.02 • 0.03 + 0.03 ~ 0.03

9.15E-5 1.05E-4 1.14E-4 1.22E-4 1.28E-4 8.43E-5 9.95E-5 1.07E-4 1.13E-4 1.21e-4 9.02E-5 1.04E-4 1.12E-4 1.18E-4 1.25E-4 9.03E-5 1.04E-4 1.12E-4 1.20E-4 1.27E-4 - - - - - - - - - - - - _

8.91E-5 1.03E-4 1.11E-4 1.18E-4 1.25E-4 • 0.32 • 0.02 + 0.03 • 0.04 • 0.03

III I I I III

TMA sample pieces for the determination of the 1.e.c. in the X-, Y- and Z-direction were machined out of these A, B, C and D samples. The l.e.c, was measured between -20~ and 120~ subsequently, after a 20~176176 thermal pretreatment to remove frozen-in stresses. The results of these measurements are listed in Table 9.9.

A repeatability of • 0.03 was measured for the 1.e.c. value determination of PK terpolymer between 20~ and 80~ This repeatability value is the reason that the 1.e.c. values in the X- and Y-direction are considered to be equal. The 1.e.c. values measured in the Z-direction proved to be about 6 % ~ower than the values measured in the X-, Y-directions.

327

9.5 Determination of electrical properties of PK terpolymers

9.5.1 The influence of moisture on the dielectric properties _

The relative high dielectric constant of water makes the dielectric constant of a polymer very sensitive for small amounts of absorbted moisture, see 5.2.1. PK terpolymer absorbs under room temperature conditions 0.5 - 0.7 %wt. of moisture. The influence of such a moisture concentration on the dielectric constant and the dielectric losses were determined.

A two millimeter thick injection moulded disk sample was, placed in a three terminal guarded sample holder and provided with golden electrodes applied by a vacuum evaporation process, connected with the automated dielectric measuring system described in chapter 5.1.4. The dielectric properties of the moisture saturated sample were measured from -100"C to 150~ using a heating rate of 0.5~ The investigated sample is nearly completely dried during this heating procedure. Besides, amorphous phase ageing effects (see 9.3.2) are also removed by this thermal treatment. A second measuring scan from -100~ to 150~ gives then the dielectric properties of the dry, non-aged sample. The weight loss of the dielectric sample, determined immediately after this second scan, was considered to be caused by evaporated moisture. A third heating scan was performed, subsequently, to measure the specific volume resistivity between -40~ and 140~ A small mass loss effect (0.07 %wt.) was still measured after this third heating scan; the moisture content of the sample was considered to be zero after this third scan. Both dielectric measuring scans were performed at nine different frequencies between I00 Hz. and 1 MHz.

The dielectric constant/temperature curves measured at a frequency of I000 Hz. are shown in Figure 9.15, nummerical values are listed in Table 9.10. A dielectric constant increasing effect due to the absorbed moisture phase is clearly present. The extent of this effect is smaller in the glass-rubber transition region (0~ - 40~ than at temperatures < 0~ and > 40~ The reason for this difference is the ageing effect of the amorphous phase during the first scan which decreases the extent of the step-wise dielectric constant increase. This ageing effect was already detected during the DMA experiments (see 9.3.2).

The dielectric loss curves (Figure 9.16) show, just like the dynamic mechanical loss curves two relaxation effects- the glass-rubber (E) relaxation effect at about 20~ and a low temperature (7) relaxation effect. Both loss curves in Figure 9.16 illustrate again that ageing decreases the strength of the E relaxation without influencing the temperature location of the relaxation effect. This ageing effect is strong enough to make the dielectric loss values of the dry, non-aged sample for temperatures above Tg higher than that of the 'wet', aged sample. At temperatures below Tg, however, the dielectric losses of the dried sample are ~ than that of the 'wet'

7.5

u~ c- O 0

0 r

O tD

6.5

5.5

4.5 -50

first/second heating scan

Figure 9.15 Influence of moisture on the dielectric constant of PK terpolymer o

/ / o

O ~ O ~"

o

0 mo I s l ; u r e ~ /

c o n t e n t - 0o

~. ~t 0 / 0 / m o ! s t u r ' e / 0 / e ~ c ~ - 0.07 Y. wt.

0 / 0 measur'tng ~Pequency: L000 Hz " / O �9 I ' I ' " 1 ' " I " ' I

-30 - 10 10 30 50 Temperature (~

O .

O / 0 / /

o' 7o / / / o

~ @

U / pO" �9

/ o /

�9 I ' I

70 90

o 0

._e o

1 0 - t -

- ? 8

Figure 9.16 o Influence of moisture on the dielectric losses of PK terpolymer e /

0 ~ - 0 ~ 0 m o t = t u r e c o n t e n t - 0 . 4 7 ~ u t .

.O f { ~ , " ~ O ~ O~

/ �9 �9 \ / . , , , , o /.o,=,~ ooo,oo, \ \ / "

4o--o\\ ," I 0 0 I ~ I " 0

\, Xo \o-~/ e

1.8 kHz. ' I " I ~ I ' ~ ' I ' ' I '

-58 -38 -10 18 38 58 Temperature (~

I ?8

330

sample.

The dielectric losses at temperatures below Tg are shifted to a higher level due to the moisture absorption. Remarkable, however, is that the intensity of the dielectrically measured 7-relaxation as such is practically not influenced by the samples' moisture content. Dynamic mechanically, a strong influence of the moisture content on the intensity of the 7- relaxation was measured, see Figure 9.9.

Table 9.10 Dielectric properties of PK terpolymer between - 20~ and 120~ at a constant frequency of I000 Hz.

I I

temperature

oC ._

- 20 m

0 �9 =

20 [] . . . . . . .

40 l l , ,

60 i

80 �9

i 0 0 ==

I I I �9 I I

die1. die1. loss constant, factor,

El E l l , ,,

5.14(4.92) 0.114(0.094)

5.55(5.32) 0.086 (0.083) , , ,

6.37 (6.26) 0.106 (0.120) , ,

6.83 (6.72) 0. 099 (0. 115) ,,, , ,, ,

7.33(6.97) 0.236 (0.308) . . . . . . . . . . . . . .

7.86 (7.41} 0.462 (0. 880) ,, , , , , ,

8.35(8.20) 0.850(1.712}

120 9.15 1.768 f I rl I I I I II

5.14, moisture content = 0.47 %wt. (4.92}, moisture content = 0.07 %wt.

,i I '

Tan ~,

E I t / E I . . . . .

0.0222 (0.0191)

0.0155(0.0156)

0.0166(0.0192) , ,

0.0145(0.0171)

0.0322 (0.0442) , ,

0 . 0 5 8 8 ( 0 . 1 1 8 6 )

0 . 1 0 1 8 ( 0 . 2 0 8 8 ) . . . . . . . .

0.1932 I I I I

331

9.5.2 The freuuencv dependency of the dielectric properties The presence of relaxation effects makes the dielectric properties not only temperature but also frequency dependent. The data listed in Table 9.11 show the extent of this effect at room temperature. The curves in Figure 9.17 give an impression of the temperature and frequency dependency of the dielectric losses of PK terpolymer due to the E- and Y- relaxation effects.

Table 9.11 Dielectric properties of PK terpolymer between 0.1 - 1000 kHz. and at a temperature of 20~ (moisture content = 0.07 %wt. )

0.3 ,. ,,

1 . 0 []

3 . 0

10 L

30 l ,,

100 [ ] . .

300

L ! .o0 .0

i i

freq. die1. kHz. constant,

E I . . . . . . .

0.1 6.26 , ,,, , | , , ,, , , , ,,

6.31

6.26 , ..

6.20

6.11 , , , ........

6.02

5.88 . , ..

5.72

5.54

I'I '

die1. loss factor

Eft _

0.209

0.143 . ,,,

0.120 .. ,

0.120 ,,..

0.135 ,

0.163

0.213 L, ,

0.277

0.355 Inl n'l rl I

Tan

s t /61

0.0333

0.0227

0.0191

0.0194 ,

0.0221 , L . ,

0.0271 , ..,.. f-

, ,,

0 048 i 0 0641 |

I I I " ~ l l l I ~ ' I

These curves show a rather complicated relaxation behaviour: - the E-relaxation seems nearly disappeared at low frequencies

due to a high level of background losses, see the 0.1 kHz. curve,

- the ~- and y-relaxation effects overlap in the 30 kHz. - 100 kHz. frequency region, see the 100 kHz. curve,

- the Y- and E-relaxation seem to behave like one single effect for frequencies > 300 kHz., see I000 kHz. curve.

The dielectrically measured relaxation maxima are plotted in an Arrhenius plot together with the DMA data (9.3.2) and the result of the Tg(midpoint) determination by DSC (9.3.3), Figure 8.36. The g-relaxation hardly shifts as a function of the measuring frequency due to the presence of a large crystalline phase. The y-relaxation is clearymuch stronger frequency dependent. An activation energy value of 63 kJ/mole was calculated for the y-relaxation from the slope of this curve. Some typical y-relaxation activation energy values for linear polymers are 63, 54 and 54 kJ/mole for respectively PVC [14], PC and PET [15]. The mechanisms of these y-relaxations are often described as local mode relaxation effects [15]. The same mechanism might also be responsible for the y-relaxation effect in polyketone polymers.

18 e- , second heating scan 0 I k Hz / v t 0 k Hz -i �9 �9 0 �9 �9

I

. / /. 0 v

~ / o . , & ~ & / o 18.0 kHz.

i ~ / .o~_ . / I .o~o-O-~ 8~I /

'" 0 " ~ / ' - - ' - I x 0 v / -, V ~ / IS -z'- / /o ~ , / N. / " ~ I ~ \ ~ o

- _v / / -"/ ~ v-- ~" \ ~" - , . I , o o / . / \ v ~ ' \ a <>. o / \, , " / " / / ~ ~ o ~ " \ <>

~ " 0 % , \ / 0 ~/ 0 ~..<>/ 180.0 kHz.

a 0 , . . , / O O ~ e , 8 0 8 . 8 kHz /

B A / Figure g. 17 Dielectric loss / temperature - frequency relations of PK terpolymer

mots t , u r e c o n t . e n r - 8 . 8 P ~ wt,.

10-Z-

-188 -88 -68 -48 -28 0 28 48 68

Temperature (~ 80 100 128 140

333

Figure 9.18 Frequency/temperature relation of the beta and gamma relaxations of PK terpolymer

+ & �9 diel, loss mech. loss DSC midpoint

1 6

1 4

12

10

"T 8 - i,,,,,,,,,,=,l

E '--' 6 - c _J

4

2

0 -

- 2 3 .20

+

+

+

+ I + I + I +

beta re laxat ion

+

X gamma X relaxation

+

N +

& I

0 I ,

3.60 | I J i

4.00 4 .40

1000/'r(max.), K -1

X \

\

|

4 . 8 0

\ \

\

\

I | ,, ! ,

5.20 5.6O

334

9.5.3 The volume resistivity determination of PK terp_ol_vmer The specific volume resistivity was measured between -40~ and 140~ during a third, step-wise heating scan. The sample/ sample cell combination used for the dielectric measurements was connected for this purpose with the automated resistivity measuring system described in 5.1.4.

The measurements, at nine discrete temperatures, were performed with an electrification voltage of 500 Volt and an electrification time of thirty minutes. The measured currents were used to calculate the so-called p-60 value and the p-dc value, see 5.1.5.

The results of the specific volume resistivity measurements on PK terpolymer are listed in Table 9.12 and plotted as a function of the reciprocal, absolute temperature in Figure 9.19. The p-dc value of PK terpolymer decreases about five decades (from about IEI40hm.m to IE9 Ohm.m) due to the change in the amorphous phase from a glassy into a rubbery state. The measured resistivity level is too high to be responsible for the low frequency 'background' losses as measured in the 0.i kHz. curve of Figure 9.17.

Table 9.12 The specific volume resistivity of PK terpolymer. (average moisture content 0.07 %wt. )

I I

temp., oC

-38

-21

- 1

19

39

63

88 .

112

136

II I

p-60, Ohm.m

3.4E14

1.4E14 , ,

9.3E12 ,,

4.9EII

i. 9El0

1.5E 9 . . . . . ,

4.1E 8 .J

3.1E 8 �9 .

2.4E 8 II I

I

p-dc, Ohm.m

6.2E15 ,,, , ,

8.5E14

6.3E13 ,,, ,,

5.6EII

2.0El0

2.1E 9 ..

1.1E 9 . , ,

1.0E 9 I

7.1E 8 j I II

The dielectric loss factor measured is the sum of a dielectric contribution and a dc conductance contribution, see 5.1.3. The dc conductance contribution to the loss factor given in equation 5.19 can be written as:

Er"(dc) - I/(w x p-dc x 8.85E-12) 9.5

335

Figure 9.19 PK terpolymer specific volume resistivity absolute temperature relation (dried, compression moulded sample)

+ A R h o ( 6 0 ) R h o ( d c )

16

15

.E 14 E e -

O

._~ m ccj i1) ft. - 1 2 0

d CI.

0 ._J

10

9

2 . 4 0

�9 I i . ! . J ............ I

2 . 8 0 3 .20 3 .60 4 . 0 0 4 . 4 0

lO00/T(max.), K -1

336

The p-dc value of 5.6Ell (Table 9.12) is used to calculate the dc loss contribution at 20oC. Thus equation 9.5 reduces to:

~r" (dc) = i/(31.12 x f) 9.10

where f - the frequency, Hz.

The measured ac/dc dielectric loss values at 20~ and respectively 100, 300 and 1000 Hz. (Table 9.11) and the calculated dc loss effects are:

frequency, Hz. or" (ac/dc) cr" (dc)

i00 0.209 0.00032 300 0.143 0.00011

1000 0.120 0.00003

Hence, the dc contribution to the dielectric loss factor at 20~ can be neglected. First at lower frequencies and higher temperatures a still relative small dc conduction effect is calculated, for example at 100 Hz:

40~ Er"(ac/dc) -- 0.454 ~r"(dc) = 0.009 65~ ,, = 1.70 ,, = 0.086 88"C ,, . 3.87 ,, = 0.164

It will be clear that the dielectric losses of PK terpolymer systems are mainly stemming from the strong dipolar relaxation effect of the carbonyl groups present in the polymer mainchain.

337

9.6 Survey of PK terpolymer thermal analytical characterisation results

A. The crystalline phase XRD and TA analysis-have both shown the semi-crystalline character of the in this chapter described PK terpolymer system. The main properties of the crystalline phase measured on non-stabilised, powdery reactor samples, are:

Tin-value (DSC, 20~ first heating) = 220 • 3~

Hf-value (DSC, 20~ first heating) = 111 _+ 10 J/g

Hf (max.) -value (calculated value) - 20~ J/g

crystallinity x (c) - 0.54 _+ 0.05

BLThe amorohous phase _

PK terpolymer shows three relaxation effects in the temperature region from -100~ to 180~ (DMA, sheet samples compression moulded at 240oC): - the u-transition, a crystalline phase relaxation effect with

tan delta maxima at temperatures between 130 to 145~ - the fi-transition, the glass-rubber transition of the

amorphous phase with tan delta maxima at about 19~ and, - the 7-transition, a secondary amorphous phase relaxation

with tan delta maxima at about -75oC.

Tg (onset) -value (DSC, 20oC/minute - second heating) = 4 _+ 3~

The E-transition is influenced both by an ageing and a moisture absorption effect: Ageing; a densification process of the amorphous phase, the strength of the g-transition decreases (tan delta (max.) from 0.08 to 0.05). Moisture absorption: plasticising the amorphous phase i.e. the temperature location of the E-transition decreases. e.w.s.* = 0.74 %wt. (due to storage at 20~ and 50 % RH) tan delta(max.) decreases from 19oC to 5~ e.w.s.* = 2.33 %wt. (immersion in ion-free water) tan delta(max.) decreases from 19~ to -8~

Both the strength and the temperature location of the 7- transition proved to be sensitive for absorbed moisture.

C. PK terp_ol_vmer sgecific m~terial properties at 20~ Lin. expansion coeff. X/Y'direction--- 1 10E-4 • 0'02

Z-direction = 1.03E-4 _+ 0.02

spec. volume resitivity = 5.6Ell Ohm.m

dielectric constant , frequency = 6.26 dielectric loss factor, is = 0.120 dielectric tan delta , 1000 Hz. - 0.0192

(*e.w.s. : equilibrium water saturation)

338

References

1. E. Drent, European Patent 121,96 (Shell), 1984. 2. E. Drent et al., J. Organomet. Chem., 417, (1991. p. 235. 3. J.A.M. Broekhoven and R.L. Wife, European Patent 257,663

(Shell), 1987. 4. P.J. Flory: Principles of Polymer Chemistry, New York,

(1953) 5. S.Z.D. Cheng- Polymer Analysis and Characterisation,

Applied Polymer Symposium 43, editor H.G. Barth, New York (1989) p. 336.

6. J.N. Hay, J. Pol. Sc.: Pol. Chem. ed., 14, (1976), p. 2845 - 2852.

7. B.J. Lonlnerts et al., J. of Pol. Sc. : Part B: Polymer Physics, Vol. 31, p.1319-1330 (1993).

8. R.C. Allen, Internal Shell Report, Westhollow Research Centre, (1986).

9. Y. Chatani et al., J. Pol. Sc., 55, (1961), p. 811. I0. E.A. Klop et al., J. Pol. Sc.: Part B: Polymer Physics,

Vol. 33, (1995), p. 315- 326. 11. L.C.E. Struik, Plastics and Rubber Processing and

Applications, 2, (1982), p.41 - 50. 12. B. Wunderlich: Thermal Analysis, Academic Press Inc., New

York (1990). 13. J. Grebowicz, Pol. for Advanced Techn., 3, (1992), p. 51-

59. 14. C.C. Ku and R. Liepins: Electrical Properties of Polymers,

Hanser Publishers New York (1987), p. 95. 15. N.G. McCrum, B.E. Read and G. Williams" Anelastic and

Dielectric Effects in Polymeric Solids, J. Wiley London (1967).

THERMO-ANALYTICAL CASE STUDIES

CHAPTER 10

339

CHAPTER I0: THERMO ANALYTICAL CASE STUDIES

I0.1 Introduction

The thermal analysis techniques treated in this book are used to study the physical properties of polymeric systems in relation to their chemical structure. These thermal analysis techniques are however also well suited for short case studies. Many of such case studies are performed in several technical service laboratories all over the world. The results of such (short) case studies are rarely reported although they often contain interesting information.

A number of such case studies, initiated by technical support questions, are reported in this chapter to illustrate the often interesting product information obtained with a limited number of TA experiments performed just to answer a 'simple' question.

10.2 The effect of the presence of a solvent during the cure of a thermoharding system.

The physical properties of cured resin systems are usually tested using rectangular casting samples and/or film samples on a metal background. The casting test samples are obtained after the cure of a proper resin/curing agent mixture without any solvent present. To make a film test sample however nearly always a solvent has to be added. The physical properties, like the Tg-value, of such film samples after cure seem often to be lower than those of the solvent-free prepared casting samples.

This effect might be caused by a difference in the network structures formed but it might also be a simple residual solvent effect. A number of DSC Tg-value determinations were performed to obtain more information about this subject.

The diglycidyl ether of bisphenol A (DGEBA, EPIKOTE 880 ex- SHELL) resin was cured with a stoichiometric amount of diaminodiphenylmethane (DDM). A thick casting sample was of about i millimeter prepared from a part of this mixture (sample A). Subsequently, about 50 %wt. of methylisobutyl- keton (MIBK; boiling temperature = I16oc) was added to the remaining part of the DGEBA/DDM mixture. A sample of one millimeter (sample B) thick and a sample of 13 micron thick (sample C, on a metal background) were made from the DGEBA/ DDM/MIBK mixture.

The three samples were cured then using the more or less 'standard' temperature treatments i.e. 1 hour/80~ hour/150~ hour/175~ for both castings and 2 hours/120oc for the film sample.

340

DSC measurements were used subsequently to measure the Tg (onset)-values of these systems (measured during the heating mode, scanning rate 20~ �9

Sample A, 1 mm thick/prepared without MIBK �9 172~ Sample B, 1 mm thick/prepared with MIBK - 70~ Sample C, 13 micron thick/prepared with MIBK: 155~

Non-isothermal TGA experiments showed that sample B still contained 13.5 %wt. MIBK while sample C still contained 4 %wt. MIBK. The TGA experiments also showed that a temperature of about 250~ is necessary to remove the residual solvent from the samples B and C within a reasonable amount of time (I to 2 hour).

Subsequently, the three samples were stored in an oven at 250oC (in a nitrogen atmosphere). The weight loss of the samples and the increasing Tg(onset)-values were followed as a function of time.

The Tg(onset)-values of the samples B and C did not further increase from the moment that the sample weights became constant. The following end-values were measured:

Sample A, wt. loss after 150 min./250~ 0.4 %wt...Tg = 172~ Sample B, wt. loss after 150 min./250~ 13.7 %wt...Tg = 166~ Sample C, wt. loss after 90 min./250~ 4.0 %wt...Tg = 165~

These experiments clearly show that the presence of residual solvent (after the standard cure procedure) is the main reason for the sometimes considerable differences in properties. The time/temperature necessary to remove the residual solvent is determined by the sample dimensions and the viscosity of the resin system. The boiling temperature of the residual solvent seems to be less important.

Figure i0.i shows the relation between the Tg(onset)-values and the amount of residual MIBK for the samples B and C. Both systems arrive, after evaporation of the solvent phase, at equal Tg(onset)-values. Figure I0.I shows, however, that these equal end-values are reached along different ways.

The Tg(onset)-value of both in the presence of solvent cured systems is indeed lower than that of the solvent-free cured system. This difference of 6~ • I~ was also measured for a second identical series of DGEBA/DDM samples prepared with bis(2-ethoxyethyl) ether (BEE, boiling temperature about 150~ as solvent phase.

We expect that the impact of this difference on the physical properties will be small because it is mainly caused by a difference in the shape of the glass-rubber transition effect. The temperatures at which the Cp-curve really starts to deviate from the base curve i.e. the real starting temperature of the glass-rubber transition effect, appears to be practically identical for all three systems!

341

Figure 10.1 The Tg-value of DGEBA/DDM resins as a function of the residual solvent (MIBK) content

+ 1 mm A 13 u o c a s t i n g f i lm

ref. cas t i ng

1 8 0

1 7 0

1 6 0

1 5 0

1 4 0

1 3 0

1 2 0

1 1 0

. . . . . .

0 0

w

r - O

I - - -

\ +

+

4-

. . . . . , . . . . . . . . . . . . . I , I -~' I , ~ j . . . . . . I ,

0 1 2 3 4 Residual MIBK, %wt.

+

5

342

10.3 The thermal transitions of a liquid crystalline polymer.

Thermotropic liquid crystalline polymers (LCP) show during heating one or more mesophase transition effects before they change after an endothermic fusion maximum into an isotropic melt. These transition effects are usually indicated by Tm, Tm' etc. for the mesophase transitions and by Ti for the transition to an isotropic melt. The presence of such a mesophase (for mainchain LCP systems usually a nematic one) offers the possibility to use these polymers as a reinforcing fibre phase in so-called 'self-reinforcing' composite systems.

The LCP system Vectra B950 (an ~romatic copolyesteramide) from Hoechst-Celanese was intended to be used as such a reinforcing fibre phase in a polypropylene (PP) matrix. We were requested to determine the thermal transition effects of Vectra B950 in order to use these effects to determine small amounts of Vectra B950 in a PP matrix by DSC.

A DSC heating/cooling scan (scan rate 20~ from 20~ to 450~ and back to 20~ with a Vectra B950 nib sample showed a mesophase transition effect and an isotropic melt effect. Two mesophase transition effects were detected in the extruded Vectra B950 systems. All three endothermic effects mentioned, see Figure 10.2, proved to be thermally reversible-

Tm' -value : 161~ Hf (m') -value: 4 J/g Tc(m')-value: 112~

Tm-value �9 280~ Hf (m) -value : 2 J/g Tc(m)-value : 227~

Ti-value �9 396~ Hf (i) -value : 84 J/g Tc(i)-value : 374~

The DSC technique was not sensitive enough to determine a reproducible Tg(onset)-value of this system. The Tg determination was performed therefore with a Polymer Laboratories DMA, see Figure 10.3. The glass-rubber transition region of Vectra B950 starts (measured at a frequency of 10 Hz. and a heating rate of 2~ at about 110~ and ends at about 170~ The Tg-value i.e. tan delta(max.) is 140~ Thus, the Tm' transition effect and the glass-rubber transition occur in the same temperature region.

The dynamic stiffness (E')/temperature curve in Figure 10.3 shows a strong decrease for temperatures higher than the mesophase transition Tm. This stiffness decrease is sufficient to allow extrusion of this polymer at 285~ Extrusion of a small amount of such a LCP with PP results in LCP particles with a fibre shape in the PP matrix.

343

5 0 0 -

4.50

4.00

~ 3.50

i 3.00 - - -

2.50 . . . . . ~ . j . ~

-r- Z.~o:

I ' 00 t 0.50

0 . t~ !

200.0 250 0

C U R V E fl ( I , H(RT ING )

. . . . . . . . . . . . C U R V E B ' ( 2 , COOL INGI

\

I i i - - :300 0 350.0 I00 0 1.50 .0


Vectra B950 nib sample

500 0

o.e5

0 40

03'3

A

v ~ : o a s o LL '*" 0 2 0 .

I 0 15

0 10

oi 0

5O0

Tin" - I l l iC Tm - 2BIBtC

(F' iRgT) H[RTING .~CriN

~ . ~ ..~ / ...,...,,.~..--. �9 ...--. e....~........

"~'-.~.,~ / .

\ / ' ~ " . . . . . . . . . . . . . . ,_ .~ / -~ .~.. i �9 ! f t , . , /

i ! T~ l (SECOND) COOL l NG SCRN

Te" - 112~C

I I I " l I ! ' I I t 100 0 150 0 200 0 250 0 300 0


Results of DSC heating/cooling scans on a Vectra B950 extruded string sample

Figure 10.2

tpll"

N"

BiD"

"4 ~U :3 Q. �9

m

t,Z"

(0o) mrqmeawel sac oc ! as t , o~

�9 i " " " " I " " ...... I

"9"

§

§

§

§ § § § § §

q=

41-

§

++ § ++

eldWes (0o06~ / u!w cj) pePlnOW uo!sseJdLuoo 0cj68 eJ],oeA e uo slueweJnseew V~E] jo s~lnseEI

80L eJnSL-I

4w

f q.

4. § § §

/

4"

,4.

4.

§

8"8

i

!2"/.

8"/. .=

8

t,'8 r- - o

. m

8"8 S~

~'S

8"S

81

345

The Ti and Tm' transitions are not suited to detect such a relative small LCP phase in a PP matrix by DSC. The PP matrix is thermally not stable enough at the Ti temperature of 396~ while the LCP Tm' temperature of 161~ coincides with the fusion region of the PP matrix. The relative weak Tm transition (280"C) is thus the only possibility to detect this LCP phase in a PP matrix by DSC.

The results of the first measurements on a PP/Vectra B950 (90/10) blend were disappointing, the fusion effect of Tm was not detectable. Subsequently we tried to increase the strength of the Tm transition by a proper annealing procedure. A heat treatement of the DSC sample of two hours at 260~ increased the Hf(m)-value from 2 J/g to about 6 J/g. This Hf(m)-value increase proved to be sufficient to make a Vectra B950 phase of I0 %wt. detectable in a PP matrix using a single DSC scan.

The Tm-value itself also proved to increase due to this annealing procedure. A series of additional experiments was performed to investigate the extent of the changes in the strength and the temperature location of this mesophase transition, see Figure 10.4.

A kind of an equilibrium situation seemed to be reached first after 1440 hours annealing at 260~ The Tm-value then increased from 280~ to 3580C and the Hf(m)-value increased from 2 J/g to 17 J/g, see Figure 10.5. These results clearly illustrated the metastable character of the Tm transition but did not contribute to a further improvement of an DSC detection method of this LCP system in a PP matrix due to the possible long sample preparation time required.

10.4 The optimal crystallisation temperature of diphenylol methane

The 'raw' diphenylolmethane (DPM) is a mixture of p,p-DPM, o,o-DPM, o,p-DPM and tri/tetramers. An HPLC analysis of such a DPM sample determined 31%wt. p,p-DPM, 40 %wt. o,p-DPM, 16 %wt. o,o-DPM and about i0 %wt. trimers to be present. Such a mixture crystallises very slowly (several days) during storage at 20~ We investigated if this crystallisation process could be accelerated to improve the process efficiency, by optimising the crystallisation conditions.

A purified DPM reference sample (p,p-DPM > 80 %wt.) showed two endothermic fusion effects (Tml = 158~ = 124 J/g, Tm2 = 108~ = 14 J/g) and a clear recrystallisation effect (Tc = 120~ = 88 J/g) during heating and cooling scans at 20~ in the DSC. The difference between the Hfl-value (124 J/g) and the Hc-value (88 J/g) shows already that recrystallisation from the melt for this relative pure sample clearly occurs slower than recrystallisation from solution. The relative weak Tm2 effect has completely disappeared during the cooling scan in the DSC.

346

1.30

I. 20

1.10

1.130

g~ O. O0

0o110

j ~ 0,70

,,~ 0 , 8 0 OP

:::p.., O. 50

O. l O

0.30

0.20

O. I0

0.00 2;

. . . . . . . . . . . . . . . - - ..... - . . . . + . . . . . . . . . . . . . . . . . . . . .

S~'g H / 2 S g C I '~ . n n e e l e d

] ~L tm ,, 34g C I ! Hr * lib J / l l t zo H - ~ + e c ~ !

l l n n l I l e a l / '~ I . . , , + o . / % i I t . . =:~e c / \ / t

. ~ , . + _ ~

i a , _ _

- / X Z H /~GI I C I In . I I II I lld 1 IIi1~ I II t m - 211G C - H e - I1 3,'11

n o t . I - h e e l e d jmmpl l Tm - 2 0 0 C - HI' " 2 J / g

+ ' : : . . . . . . . . | - : . . . . . . . . . . . . . . . - | . . . . : : - - l i . . . . . . . . . . . . . . . - " . . . . . . . . I - " ........ : . . . . . . . . . . - " -

+0. 0 275. 0 300. 0 325. 0 350, 0 375. 0

Temperature (~ Figure 10.4 Effect of annealing a Vectra B950 fibre sample at 260~

t. tO

t.O0

o9o A

IC~ 0 . 8 0

0 ? 0

0 . 6 0

�9 ~ o ~ |

:I: o . J o

0 . 3 0

0 20

0 . I 0 j . . . . . . . . . . . . "- ~ '=--+ - . . . . . . . . . . . . . . .

0 . 0 0 . I + . . . . . . _ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . - ..... _ . . . . . . . . . . . . . . . . . . . . . I I I I I I I I I

2 0 0 . 0 2 5 0 0 3 0 0 . 0 3 5 0 0 d O 0 . O

Temperature (~

Xl 346 .266 "C lt2 370.000 *C

Peak 350.03! "C Tm ,,, 3511 C 1t4 16.9416 d/g HI' ,- I? J / G

1' ! - 3911 c - I ;2 J t 1 1

J

r ]

+ _ +

I 454) o

Figure 10.5 First DSC heating scan on a Vectra B950 fibre sample after 1440 hours annealing at 260~

347

These results indicate that it would be very difficult to measure dynamically recrystallisation effects of 'raw' DPM samples in the DSC.

A 'raw' DPM sample recrystallised during storage at 20~ showed an endothermic fusion effect between about 40~ and 140~ with maxima at 125~ and 94~ and a Hf-value of 105 J/g. Recrystallisation effects during cooling scans were, as expected, not detected during these experiments.

The extent of the fusion effect after a certain storage time and temperature did give however a clear indication of the progress of the recrystallisation effect. Figure 10.6 shows the fusion curves of 'raw' DPM mixture samples which were, after heating to 160~ stored during respectively i, 16.5, 24 and 65.5 hours at 20~ The slow increase of the crystalline phase is accompanied by a decrease of the extent of the amorphous phase i.e. the strength of the glass-rubber transition has to decrease at the same time. This is confirmed by the results shown in Figure 10.7.

After heating to 160~ 'raw' DPM samples were subsequently cooled at maximum speed and stored at different temperatures while the increase of the Hf-value as a function of the storage time was followed. The results of these experiments are plotted in Figure 10.8, the curve measured at 50~ is omitted for clarity reasons.

These results clearly show that the recrystallisation speed of this DPM mixture reaches a maximum value at about 60"C. The Hfl-value of 124 J/g measured for the reference sample indicates that a Hfl-value of 48 J/g for this 'raw' DPM sample containing 31%wt. p,p-DPM might be possible. The results presented in Figure 10.8 show that after two hours storage at 60"C a Hf-value of 49 J/g was measured. This indicates that the recrystallisation process of the p,p-DPM phase is nearly completed is after these two hours of storage at 60~ Thus, a solid 'raw' DPM product can be obtained within a reasonable amount of time (about two hours) if the recrystallisation process occurs at 60~ instead of 20~

348

o ~ . . . . . . . . . . . . . i oo o t t o., ~.o' 1 .___~~ y ~ t ~' �9 o o, 1 , , , , , /~: ;~~, :~~,~--~,~~ t~.

o 1 ~ ~ / ~ y - - ~ o 0 4 " 0.4

0.~ 0.3

0.2 0.2

0 I V / ( ) s t o r n g e t i m e . t r o o m t e m p e r . t u r e t 0"1

O0 I I I I ........ 1 "0"0 70.0 00.0 90.0 100,0 110,0 120.0 130.0 140.0 150.0 160.0

Temperature (~ Figure 10.6 Fusion curves of at ambient temperature recrystallised crude DPM

0"50 i O. 45

0.40-.,

0.35 A

~ 0 . 3 0

0 0.25 . . , . .

LL

020

0.15

PERKIN-ELMER 7 Sedes Thermal Analysis System . . . . . . . .

f o.50

o. 45

( I ) r0. 40

. / __~r ]s_s.~----

( 2 4 ) -0.30

/ (ss. S T - - ~ - ~ r0,~

-0.20

O. I 0 1 I

0"05 1 ( ) s t o r a g e t i m e a t r o o m t e m p e r l t u P e

o oo ~ v i I I t -40. 0 -:30. 0 -20. 0 - tO. O O. 0 IO. O

Figure 10.7 Temperature (~

I I 20.0 :310. 0

I .O. 15

0,10

ii ~ 40. 0

Glass-rubber transition curves of at ambient temperature recrystallised crude DPM

349

Figure 10.8 Hf-value development of "raw" DPM as a function of the recrystallisation time/temperature

+ A 0 + & 80~ 70~ 60~ 40~ 30~

5 0 .... ()

. 0 j

. A ~ ~ L~

_ = _ _

+ / . ~ ~ "

, I I , I . . . . . . , I L i I ,

4 0

�9 m 30 ,.T

"~ 2 0

10

0

0 2 0 4 0 6 0 8 0 1 0 0 1 2 0

Isothermal crystallisation time, min.

350

10.5 The dynamic stiffness of ultra-high molecular weight polypropylene in its melt

Ultra-high molecular weight polypropylene (UHMW-PP, Mn about 5,000,000) is used in applications where a high melt strength is necessary. The dynamic stiffness i.e. the storage shear modulus (G'), is related with the melt strength and can thus be used to compare the properties in the melt of different UHMW-PP systems. A number of experimental UHMW-PP systems was investigated in this way.

The measurements were performed with an automated torsion pendulum (DMA) apparatus working at a frequency of about 0.5 Hz. (varying frequency system, see Chapter 4). The use of a low-stiffness suspension wire permitted stiffness measurements as low as 5E4 N/m2. These DMA measurements were performed from 50~ up to 250~ using a heating rate of l~ the samples were, during these experiments, purged with nitrogen.

The dimensional changes of the sample strips during the heating of the samples through their melting regions made calculation of the dynamic stiffness in two steps necessary. The sample dimensions at room temperature were used to calculate the storage shear modulus values between 50"C and 170~ The sample dimensions after heating to 250~ were used to calculate the storage shear modulus values between 170~ and 250oc.

Figure 10.9 shows the results of the measurements on an experimental UHMW-PP sample as such and filled with 20 %v. of mica. A standard J-grade PP, filled and non-filled, was measured as reference system. The dynamic stiffness of both reference systems is for temperatures below 160~ higher than that of both UHMW-PP systems due to a difference in crystallinity, x(c) as determined by DSC-

UHMW-PP , x(c) = 0.33) J-grade PP , x(c) = 0.45)

) based on a Hf-value of UHMW-PP/20 %v. mica , x(c) = 0.43) 188 J/g for x(c) = 1.0 J-grade PP/20 %v. mica, x(c) = 0.52)

The crystallinity is promoted by the mica addition (mica is, just like talc and carbon black acting as a nucleating agent for PP, see Chapter 1.3.2). This effect is also detected by an increase of the DSC recrystallisation temperature (Tc-value) from 108~ to 120~ for the UHMW-PP and from 106~ to 121~ for the standard J-grade system.

Both reference samples fused completely out of the sample clamps at tempertures between 160~ and 170~ The dimensions of the URMW-PP samples did change considerably (length not changed, width -23 % and thickness +40 %) but the shape of the samples was still rectangular, which permitted us to continue the measurements up to 250~ The consideable difference in dynamic stiffness (and hence in strength) between the standard

351

Figure 10.9 Dynamical stiffness of PP grades as a function of the temperature

+ PP Z~ UHMW 0 P P / 2 0 % "1' PP mica

U- - t /20% mica

l e + 0 9

= l e + 0 8

o E I , , , , ,

m

l e + 0 7 o

10000QO

1 0 0 0 0 0

~ II �9 I

~ § O ~ ~ § 0 ~+~

+ ~ +.~ + ~

- \ 2 : "+

+ " I

I

__ = I = I , ~= _ I , I .._= I

----_. +

s \ s

0

, 1 , | , I , _ ' I , l

50 90 130 170 210 I _

25O

Temperature, ~

352

Figure 10.10 Dynamical stiffness of UHMW PP/J-grade PP blends as a function of the temperature

-4'- P P A U/J . - -g . 0 U / J - g . '4" U / J - g . J - g r a d e 2 0 / 8 0 4 0 1 6 0 6 0 / 4 0

I e+081 ~

le+07

1000000

100000

10000

+

+ L

\o I ~ ~, I ~

I I I

j ... I j, t I

t I

.... I , I I

150 170 190 210 I I

23O 250 Temperature, ~

353

J-grade PP and the UHMW-PP, both in their melt, is clear. This difference even increases due to mica addition.

The nearly constant (about 5E5 N/m2) storage shear modulus level of the UHMW-PP between 170"C and 250~ is assumed to be a part of the so-called rubbery plateau of the amorphous polymer. The length (in ~ of this plateau depends on the molecular weight. The modulus level in this plateau depends on the number of chain entanglements which are acting as temporary physical crosslinks. However, also the number of these entanglements increases with the molecular weight of the polymer. Hence the shape of this rubbery plateau is completely molecular weight dependent [1,2].

DMAmeasurements on a series of UHMW-PP/J-grade PP showed that such a physical network is able to "carry" a considerable amount the standard J-grade PP before the rubbery plateau of the UHMW-PP collapses. The results in Figure i0.i0 show that for a blend of UHMW-PP/J-grade PP (60/40) the rubbery plateau is still present up to 250~ This rubbery plateau starts to disappear quickly for J-grade concentrations higher than 40 %wt.

354

10.6 The effect of an anti-static additive on the electrical resistivity of a polystyrene foam

Polystyrene (PS) foam is only accepted as packaging material for electronic components if the electrical resistivity is low enough to prevent components damage due to electro-static charges. Anti-static additives are added to PS foam batches to reduce static charge and discharge effects to an acceptable level but reliable electrical resistivity values of such foams were not available. Hence, a series of measurements was performed to determine such values.

A series of four PS foam samples with a density of 20 g/l was made using the so-called vacuum process. Neostatic HBI55 (an alkyl substituted, kationic ammonium salt) was, as the anti- static additive, added in concentrations of respectively I, 2 and 4 %wt. This anti-static additive was added, dissolved in a constant amount of water, to the expanded PS particles and thoroughly mixed. This mixture was used to press PS foam block samples of 30 x 30 x 5 cm. Subsequently, samples of 9 x 9 x 0.42 cm were cut from these PS foam blocks to measure the electrical surface and volume resistivity. These samples were first stored, however, during six weeks at a temperature of 20

1 ~ and a relative humidity of 50 %, to reach equilibrium moisture uptake conditions.

The foam samples were mounted for both surface and volume resistivity determination in a Keithely 6105 measuring cell. The electrode configuration of this cell consists of a fixed high-potential electrode and a spring-loaded low-potential electrode, both completely guarded. The samples were clamped between the high potential electrode and the guard ring of the low-potential electrode using 4.0 mm spacers. The 4.2 mm thick foam samples were thus slighly compressed in a reproducible way (between the high potential and the spring-loaded low potential electrodes) in order to obtain a good sample/ electrode contact.

All measurements were performed with an electrification voltage of i000 Volt and an electrification time of twenty minutes. This electrification time proved to be sufficient to avoid contributions of time dependent charging currents (see Chapter 5.1). The in this way performed measurements agree with the recommendations given in ASTM D257 and IEC 93 methods.

The nummerical results of these measurements are listed in Table I0.I. Figure i0.II shows resistivity values plotted as a function of the anti-static additive concentration. The addition of this additive certainly works; the surface resistivity decreases about four decades and the volume resistivity decreases about eight decades due to the addition of 4 %wt. of Neostatic HBI55 anti-static additive. The shape of these curves indicates that addition of higher concentrations of this additive will be hardly effective.

355

Figure 10.11 Resistivity of PS foam as a function of the anti-static additive content

-'t- s u r f a c e A vo lume r e s i s t i v i t y r es i s t i v i t y

l e + 14

l e + 1 3

E r

0

:=~ l e + 1 2 i=,,=,

t~

.~_ l e + 1 1

l e + 10

I e + 0 9 1

I Volume/surface resist ivi ty " 4" correlation

le+ 15 . . . . . . . . . + " I.

le'+ 14 / "

i iii \ \ l e + 0 9 ' e + l O le,'0-11 'le,"*'-12 l e . + 1 3 l e " t " ;

~ . . ~ _ Specific surface re.iStivity. ~ -- \\

I ......... I ..... I. ,, I

l e + 1 5

l e + 1 4

l e + 13

l e + 1 2

l e + 1 1

l e + l O

l e + 0 9

l e + 0 8

l e + 0 7

0 .50 1.50 2 .60 3 .50 4 .50

E t-- 0

._> t~ ,===

e= q) E

m >o 0 , m

o,=== o G) Q.

or)

Anti-static additive conc., %wt.

356

Table i0.i Results of electrical resistivity measurements on PS foam samples

Neostatic HB155

I content,0 %wt. �9

spec. volume resistivity,

Ohm. m , ,

6.6E14 9.5E 7 6.0E 7 9.4E 6

= , , , , , , ~ . . . . . .,

spec. surface resistivity,

Ohm ., ~ ..

3.4E14 2.7El0 7.2E 9 1.9E 9

Physical measurements performed on foam samples instead of solid, isotropic polymeric systems nearly always show an increasing scatter of the results with a decreasing foam density. Besides, the scatter of the surface resistivity measurements is always higher than that of the volume resistivity results due to the higher sensitivity of the surface resistivity determination for the sample/electrode contact quality. The relative good correlation found between the measured volume and surface resistivity values, see the inserted figure in Figure i0.ii, indicates that the electrode configuration used resulted in reproducible contact conditions during this series of measurements.

10.7 The dielectric constant of polyethylene foil

One of the first polymeric products in which a relative high amount of recycled product was used, was low density polyethylene (LDPE) foil. In connection with the calibration of foil thickness transducers) it was necessary to determine the influence of a certain amount of recycled product (replast) on the dielectric constant of the foil.

Accurate dielectric constant measurements on polyethylene foil samples are not easy. The apolar character of this polymer with a very high electrical resistivity and very low dielectric constant/losses requires high resolution measurements to detect the potentially small differences.

These capacitance measurements were performed therefore with a General Radio GRI621 Precision Capacitance Measurement System with a resolution better than 0.0001 pF. The accuracy is at least 0.001 pF and the temperature sensitivity is less than 0.003 pF/~ The measuring frequency of this system can be varied between i0 Hz. and I00 kHz.

The liquid displacement method, described in ASTM D1531-62, is in fact the only proper method for accurate dielectric constant measurements on the thin (about 0.2 ram) foil samples. A gold plated ERA liquid cell with an electrode spacing of 1.435 mm was used for these measurements. The measuring

357

temperature was 21 + 0.4~

The density of the foil samples was also measured in order to correlate the dielectric constant results with the density values. These density measurements were performed according to ASTM D1505-68 with the aid of a water/ethanol density gradient column at a temperature of 21.00~ _+ 0.01~

The frequency dependency of the measuring method was first investigated. The dielectric constant (E'r) of a CARLONA LDPE 25002FA (with 1.5 %wt. Ti02) was measured twice between i0 Hz. and I00 kHz.-

CARLONA LDPE 25002FA + 1.5 %wt. Ti02

~'r-value, 21~ I0 Hz.: 2.339 and 2.345

i00 Hz.: 2.332 and 2.334 1 kHz.: 2.330 and 2.332

10 kHz.: 2.329 and 2.331 100 kHz.: 2.320 and 2.322

The ~'r value of the apolar polyethylene should be frequency independent. Von Hippel [3] reports c'r values of 2.25 and 2.26 for a frequency region from I00 Hz. up to I000 MHz. An electrode polarisation effect [4] might be the reason for the relative high ~'r values measured for frequencies ~ I00 Hz. The decrease of the ~'r values for frequencies ~ I00 kHz. is probably caused by the effect of parasitic measuring cell capacities which increase at higher frequencies. The measuring results in the frequency region between I00 and I0 kHz. are therefore considered to be the most reliable values. All measurements were performed subsequently with a measuring frequency of 1 kHz.; the results are listed in Table 10.2.

Lanza and Hermann [5] found for polyethylene a linear relation between the density and the dielectric constant according to.

E'r = (2.01 x p) + 0.427 I0.I

Wurstlin investigated the validity of this relation and reported [6] that his experimental results agreed with the values calculated with equation I0.I.

The dielectric constant values of the three pure i.e. non- filled LDPE grades in Table 10.2, are calculated with equation i0.I. The agreement of these values with the experimental values is good.

Addition of Ti02 results in a density increase and in a slightly stronger increase of the dielectric constant. Both experimental dielectric constant values of the Ti02 filled samples can be calculated, however, with equation I0.i if the value of the added constant is increased from 0.427 to o.446. Subsequently, the dielectric constant values of the LDPE foil samples with different amounts of replast were calculated using this modified relationship.

2.35

"0

m

:3 t~

0 v

r

r 0 0 0

ID ._~ D

Figure 10.12 Correlation between measured and calculated dielectric constant values of LDPE

4- L D P E A L D P E / T i O 2 O L D P E / T i O 2 / fo i l ' r ep las t '

2.34

i i , i

_ 0

. 0

2.33

2.32

2.31

2.30

2.29

2.28

2.27

358

2.27 2.29 2.31 2.33 2.35

Dielectric constant (measured)

359

Table 10.2 Dielectric constant values of LDPE foils at 21~ and a frequency of 1 kHz.

LDPE foil sample

LDPE non- filled CARLONA 18003GA DOW 150

~j bOW ,L, ~'X. 2 ~

CARLONA 25002FA/TiO2 + 0.7 %wt. Ti02 + 1.5 %wt. Ti02

CARLONA 25002FA with 1.5 %wt. Ti02 and- + 0.0 %wt. replast + 4.7 %wt. replast + 12.8 %wt. replast + 22.7 %wt. replast

i+ 32 -9 %W t- replast n ,,,

p-value, g/cm3 ........

0.9210 0.9208 0.9404

,, ,,,

0.9327 0.9384

0.9386 0.9386 0.9305 0.9349 0.9326 f ~ ii i,, ~ ' : ,, ,,

C ' r-value, measured

2.280 2.279 2.318

...... ,

2.322 2.330

2.332 2.324 2.322 2.321 2.319 "~! , h ' ,',"!

' r-value, calculated

2.278* 2.278 2.317

. . . .

2.321"* 2.332 .

2.333** 2.333 2.316 2.325 2 . 3 , 2 1 . . . . . . . . . j

* calculated with constant- 0.427 ** calculated with constant: 0.446

The effect of the addition of recycled LDPE on the dielectric constant of the resulting foil materials proved to be small. The dielectric constant value decreased for the investigated system from 2.332 to 2.321 (average value); a decrease of 0.47 % due to the addition of up to 33 %wt. of recycled LDPE.

The correlation (see Figure 10.12) between the measured and calculated dielectric constant values certainly decreases indicating an increase in the scatter of both the dielectric constant as well as the density values due to the addition of a certain amount of recycled LDPE.

10.8 The volume resistivity epoxy based moulding powder systems during immersion in hot water

Retaining a high electrical volume resistivity in the presence of moisture is one of the most important properties of moulding powder resin systems which are used for the encapsulation of electronic components.

The resistivity measuring procedure to compare such moulding powder systems within an acceptable period of time is accelerated by immersion of the samples in hot water. The volume resistivity of three epoxy based moulding powder systems was measured as a function of the storage time in water at a temperature of 90"C.

(/'3

"rl

"13 --(I

) (I)

_.

.Jk

0..0

m

~

�9

(I)

CO

:3

===D

~

0 3 3

OT

0'~

I �9

I

�9

L~

0"~

O

,I~==

=,,-

==,=

~

3 ~3

361

Resistor shaped test samples with two embedded wire electrodes, see Figure 10.13, are used to perform these experiments" These test samples were prepared at a mould temperature of 180~ and a compression time of two minutes. Subsequently, the samples were post-cured during four hours at 180oC.

The volume resistivity was determined by measuring the current flowing 60 seconds after application of a i000 Volt measuring voltage i.e. the so-called p(60) -value. This value is calculated according to-

p - (v x A)/(I x d) I0.2

where: V = applied voltage, Volt I = measured current, Ampere A = electrode area, m2 d = electrode distance, m p = a specific volume resistivity

value, 0hm.m

The electrode area and distance are approximated by considering only the two opposing surfaces of the embedded wires as electrodes. Substitution of the known values for V, A and d results then in-

p - 1.31/I I0.3

The samples were immersed in demineralised water for the ageing experiments and placed in an oven at 90~ • 2~ After a certain immersion time, a sample was taken out of the oven and cooled in water at a temperature of 20~ Subsequently, the sample was surface dried, connected with the measuring equipment and measured. This procedure was performed in about five minutes. A separate sample was used for each measurement. It is important to realise that although the samples were stored in water at 90~ the resistivity measurement itself was performed at room temperature. The measuring procedure was kept as short as possible, p(60) i.e. some time dependency of the measuring currents could stil be measured, to keep the moisture evaporation losses as low as possible.

The results of the measurements on the three epoxy based moulding powder systems are listed in Table 10.3. The same results are plotted as a function of the immersion time in Figure 10.14. The shape of these curves is clearly different. The best system (A) retains a high resistivity value after more than 900 hours of immersion. The worst sample (C) clearly shows a kind of break-through effect after about 500 hours of immersion. System B is showing an intermediate behaviour.

362

Figure 10.14 Resistivity decrease of moulding systems versus the hot water contact time

+ system a system o A B

sys tem C

l e + 1 2

l e + 1 1

E , ,C

0 l e + 10

.>_ c t }

E -~ 1 e + 0 9 0

l e + 0 8

l e + 0 7

+

0~0 ~ + ~ + "~0

A

0

0 , , j l , , j , . I . j . , , I ,

2 0 0 4 0 0 6 0 0

, +

I !

8OO 1000

Immersion time, hours

363

Table 10.3 Results of the resistivity measurements on moulding powder systems after immersion in water at 90~

Immersion time,

hours

1 6 0 2 3 2 4 9 8 6 9 2 9 0 5

. . ...... ' ,,

volume resistivity system A, Ohm.m

.. ,

1.3E12 2.2Eli 3.9E10 2.0El0 7.7E 9 5.5E 9 5.5E 9 r ., r., r

, ~ I . , ,

volume resistivity system B, Ohm.m

2.6E12 1.2El0 2.8E 9 2.2E 9 1.3E 9 9.4E 8 4.9E 8

, I., , . ,,

,,, ,, J . , ....

volume resistivity system C, Ohm. m

1.9E12 8.2E 9 4.1E 9 3.7E 9 3.1E 9 3.3E 8 2.3E 7

., . .....

364

10.9 The determination of the composition of a car-tyre rubber

Carbon black filled, vulcanised rubber is still difficult to analyse with standard analysis techniques like FTIR and NMR. Application of the thermal analysis techniques TGA and DSC offers the possibility to obtain a reasonable impression of the composition of such a rubber sample. This is illustrated below by the results of TGA and DSC measurements on a rubber sample from a Michelin MXT 185/65-R14 cartyre.

A sample of about i0 milligramme was subjected to the following temperature/time programme- - heating from 20~ to 300~ at a rate of l~ and in a

nitrogen atmosphere, - thirty minutes isothermal at 300~ - heating from 300~ to 480~ at a rate of l~ - five minutes isothermal at 480~ and switchting from

nitrogen as sample purge gas to air, - heating from 480~ to 980~ at a rate of I~ The so-called extractables or oil fraction are removed from the sample due to evaporation during the first two programme steps. The polymer fraction is, subsequently, removed during the third programme step due to thermal degradation. The carbon black fraction burns completely during heating from 480~ to 980~ in an air atmosphere. The anorganic components present in the rubber form, finally, the measured residues.

The TGA experiment was performed during the night and took 15 hours and 35 minutes 'instrument' time. The results of this experiment are listed in Table 10.4. Figure 10.15 shows the DTGA curve measured during the thermal degradation of the polymer fraction (300~ - 480~ This DTGA curve shows two minima; the second minimum is slightly a-symmetrical on its low-temperature side. The temperature locations of both DTGA minima indicate that this cartyre sample is a blend of BR and NR rubber.

Subsequently, about 15 milligramme of sample was used for a DSC Tg(onset)-value determination (heating rate 20~ This Tg-value determination was used to confirm the possible presence of a BR and a NR phase.

All the values measured of this sample and the necessary reference values are collected in Table 10.4 The two separate Tg-values at -99~ and at -64~ confirm that the vulcanised rubber sample consists for 52 %wt. of a BR/NR blend (the Tg- value of the BR phase might be shifted to higher temperatures due to the oil addition; the Tg-value of the NR phase is hardly influenced by the oil addition). The sample also contains 15 %wt. of oil and 32 %wt. of carbon black. About 25 milligramme of sample and about 16 'instrument' hours were necessary for this analysis.

PERKIN-ELMER 7 Series Thermal Analysis System

-0. !

A e~

o ~

v

.>_

! . . _

o

-0, 2

-0, :8

-0 , 4

-0. 5

L~

-0. 6 ItTX tO5 /S5 -R 14

t Figure 10.15 -o. 7- DTGA/temperature curve [ . . . . . . . . . . . I . . . . . . . . . . ~ . . . . . . . I . . . . . . . . . . 1 ......

3 0 0 . 0 325 . 0 3 5 0 . 0 375 . 0 4 0 0 . 0


i 42S. 0

I ! 450. 0 475. 0

366

Table 10.4 Results of TGA/DSC analysis on a cartyre rubber

,,, ,, , ,,,, ,. ',., , , ,

1. TGA results extractables (oil phase) , %wt. polymer phase , %wt.

organic material, total , %wt.

I . . . . r

14.5 52.0

66.5

carbon black , %wt. residue , %wt.

31.5 2.0

.. anor@anic material, total , %wt..

2. TGA results DTGA minima (primary) , oc

(secundary) , oc DTGA reference minima BR rubber (97 % cis) , ~ SSBR rubber (23 %wt. styrene) , ~ NR rubber (primary minimum) , ~

(secondary minimum) , ~

i . . . . . , ,,

3. DSC results Tgl (onset) -value , ~ Tg2 (onset) -value , "C DSC reference values BR rubber (vulcanised/no oil) , ~ NR rubber (vulcanised/no oil) , ~

,. f I , I .. '

33.5

448 360

440 430 345 397

P

-99 -64 i

-108 -66

|,

10.10 The thermal stability of ASB

It is necessary sometimes, to increase the polarity of an originally non- or low-polar polymer, for example to print text on polymer films. One of the possibilities to realise a polarity increase is carboxylation of such a polymer with 3- azidosulfonylbenzoic acid (ASB), see Figure 10.16. After loosing its nitrogen atoms, the ASB residue reacts with the polymer molecules; the number of acid groups bound per polymer molecule determines the polarity of the carboxylated system. A series of such carboxylation experiments with PP prompted us to investigate the thermal stability of ASB itself and that of ASB in contact with PP.

About four milligramme of ASB was, during a non-isothermal TGA experiment, heated from 30oc to 500~ with a heating rate of 5~ Using a higher heating rate resulted in an 'explosive' degradation reaction. The mass/temperature curve showed a two-step degradation process with DTGA minima at 191~ and 320~ A mass loss of 24.4 %wt. was measured between 1400C and 220~ The loss of the three nitrogen atoms per atom ASB should results, however, in a mass loss of only 18.5 %wt. Hence, there happened more than only loss of nitrogen atoms during this process.

367

Figure 10.16 3-azidosulfonylbenzoic acid

COOH

I

O - S - O I N II N II N

0. 0OI3

A

o. oo10 -

O. O00g "

O, 0005 -

O. O001

O, 0000 ' 4OOO

l O 0 . O ,,

368

9 5 . 0

~ . 0 -

0 5 . 0

0 0 . 0

, ' 5 . 0

I I I ,,

3500 3000 leJO0

C 0 2 , 2 3 6 4 �9 O. 0013

O. 0000 I o o o I'rJO ISOO 1250 tO00 "rJO 450

Figure 10.17 cm" The average FTIR gas-phase spectrum measured between 175~ and 180oc during an A S B - 180~ experiment

- O. o o l o

5 0 2 , 1 3 7 6 , : m - i " 0.0000

J 1177 c m - t

2 1 3 2 c m - I � 9

- O. O00S

O. 0003

I ,11 I I I ! I

1z4r

5

J ~ ~ ~ ' �9 . . . . * ' = ~ i ' I

1 3 7 0 c rn "1

~ . / co2-=,,,.v.

O. OOOG _

f / ~ \ | e . e e e 4 - so2-=,...-,,. ; / ,, \ , , ' / - \ \ i l ', \ k

to t5 20 25 3O 3'3 42

Time, minutes Figure 10.18 The "I'GA mass/time and the FTIR S02-CO2 absorptio~time curves measured on ASB during a 180~ experiment

C 0 2 8bS. 8t 2364 c rn"

O. 8 e 3 2

0 . 0 0 2 4

8 , 0 g l 6

8 . 0 9 0 6

O . O 0 0 g

369

ASB samples were heated, subsequently, in a TGA - coupled - FTIR system, to respectively 160~ 170~ 180~ 190~ and 200~ The heating rate was 5~ T(isothermal) was maintained for about twenty minutes until no further mass losses were detected.

Figure 10.17 shows the FTIR gas-phase spectrum measured during the 180~ experiment. Clearly present in this this spectrum are absorptions due to C02 (2364 cm-l), S02 (1376 cm-l) and an azido effect (2132 cm-l). The presence of C02 is pointing at decarboxylation of the acid group. The presence of S02 is indicating that at 180oc already a further thermal degradation of the ASB molecules occurs. The azido absorption (2132 cm-l) and the absorptions at 1765, 1348 en 1177 cm-I indicate that a part of the ASB spontaneously evaporates.

Figure 10.18 shows the TGAmass/time curve during the 180oc experiment together with the intensity/time curves of the S02 and C02 absorptions. It is clear that the mass loss due to nitrogen loss is starting before the C02 and S02 production starts. The S02 was less than ten minutes detectable in the TGA purge gas while the decarboxylation process took about twenty minutes. These two effects proved to be nearly equal for all experiments performed between 160~ and 200~

Integration of the absorption intensity/time curves and using calibration curves permitted to determine quantitively the amounts of C02 and S02 released during these experiments-

experiment ) 160~ 1.0 %wt. SO2 and 1.4 %wt. C02 with ) 170~ 1.2 %wt. S02 and 1.8 %wt. C02 T(isothermal) ) i80~ 1.0 %wt. S02 and 2.3 %wt. C02

) 190~ I.i %wt. S02 and 3.5 %wt. C02 ) 200~ I.i %wt. S02 and 3.5 %wt. C02

The S02 production proved to be more or less constant over the temperature region investigated; it is the result of the decomposition of about 3.5 %wt. of ASB. The acid group (-COOH) decarboxylation process is clearly temperature dependent. The amount of C02 released during the 160~ experiment is the result of the decarboxylation of 7.2 %wt. of the ASB sample. This amount increases to about 18 %wt. of the ASB sample during the 190~ experiments.

Subsequently, these measurements were performed on PP/ASB mixtures. A PP powder samples was mixed with i0 %wt. of ASB on a rollerbank. About 40 milligramme of this mixture was used for the TGA experiment i.e. the total amount of ASB present (4 milligramme) was the same as used during the first series of experiments.

Figure 10.19 shows again one of the FTIR gas-phase spectra measured during the 180oc experiment. Both the C02 and the azido absorptions are clearly present in this spectrum. The S02 absorption, however, is now clearly not present. During all experiments was shown that no S02 was released if ASB

370

0 .00 |6 "' . ' . , ' ' - '|= u = ' i �9 i " ' u " i

A r 2364 om- I

0 . 0013

! 176 c m - t

m - |

0.0003

' 0 . 0056

0 .00 t3

O.OOlO

0 . 0 ~ 1

O . m 3

0 .0000 0 .0000 4000 3000 ] 000 ~ WOO 17110 51900 1160 t 040 ?IS0 480

Figure 10.19 cm" The average FTIR gas-phase spectrum measured between 175~ and 180~ during a PP/ASB 9 0 / 1 0 - 180~ ex~dment

371

decomposes in close contact with PP. The excess of reaction possibilities with polymer molecules seems to prevent the ASB decomposition process.

The decarboxylation process of the acid group however proved to be not influenced be the presence of a polymer phase. These experiments thus clearly showed that especially the thermal instability of the acid group (decarboxylation) influences the efficiency of the PP carboxylation process.

i0.ii The thermo-analytical characterisation of a maize based, ' green' polymer

There is a growing interest for biologically degradable polymer systems due to a combination of increasing environmental conservation demands and the continuously increasing technological progress of such so-called green polymers. Several research groups in the United States, Japan ans Europe are studying the production processes and properties of these green polymers. A maize-based polymer film sample, made in Japan, was investigated thermo-analytically in order to follow these developments.

The film sample investigated proved to be a semi-crystalline polymer system with a DSC Tg(onset)-value of 16~ The DSC Tm(1)-value was 146~ with a Hf(1)-value of 18 J/g. The so- called 'processing-window' of this polymer proved to be about 50~ from 150~ to about 200~ This bio-polymer proved to be (as usual) strongly moisture dependent. The equilibrium water saturation of the film in contact with water was 64 %wt., however visual detectable swellings effects were not noticed.

A TGA - coupled - FTIR/MS experiment was subsequently performed to get an impression of the thermal stability and the chemical structure of this material. A starch sample (p.a., ex-Merck) was used as reference material.

The TGA results in Figure 10.20 show a mass loss of 5.8 %wt. due to the loss of moisture absorbed. Besides, the DTGA curve shows two weak effects at respectively 200oc and 236~ which were not detected for the reference sample (the degradation onset temperature of the reference sample was about 250oC).

The TGA purge gas mass~time i.e. temperature curves in the Figures 10.21 A and C confirm the loss of absorbed moisture between 30~ and 100~ (mass intensity maxima of m/z = 17 and 18 during the first fifteen minutes). The Figures 10.21A and 10.21B show at 200~ (35 minutes) a clear maximum for m/z = 17 i.e. a NH3 concentration maximum and at 230oC (40 minutes) a clear maximum for m/z = 44 i.e. a C02 concentration maximum. The MS results in the Figures 10.21B, 10.21C and 10.21D also show the release of C02 , H20 and some CO during the main thermal degradation process (60 minutes/330oC) of this on maize based polymer sample.

100 .0 -

g 5 . 0 "

9 0 . 0 -

85. 0 -

80. 0 -

7 5 . 0 -

7 0 . 0 -

6 5 . 0 -

6 0 . 0 -

t -

~ (9

/ /' \ / \ / \ /

38, t 65 ~ 1"1 ,C

T2 131. 686 C gg. g79 Wt;. Z

Y2 ~ g4. 175 Wt.. Z

t I I I I I I 5 0 . 0 100 .0 150. 0 200. 0 250. 0 300, 0 350. 0

A r

.>_.

=> Q

I I

I

Temperature (~ RATE +: s.o'cJm~ Figure 10.20 The mass loss (TGA) and mass loss rate (DTGA) of the maize-based polymer sample during heating in a helium atmosphere

- 0 . 4

" 0 , 2

" 0 , 0

- - 0 , 2

- - 0 . 4

" -0 . 6

-0 . B

- - 1 . 0

L~ ,,,.3 b3

373

e, T i t O. II~IB o . m (A) = 17, U.E. NH3

O . 4 N ~

l . I - - - - . . . . :.- " + . , J, + + ,- ,_-,- . . . . ~ " - , - II III Ill ~ Sll II 7 1

+ _ TOTA+~ t l ~ . . = "1~ ._ , , ml.~. , ,

llt~Inl e+_mm+T r .UC .: O, CPe CYCU[ ". e,W ....

t , I

I ~ l

I . I i

I ~ , PlMM+CX: I 1 r I

l . I

I. l ' I j l l l

I . I |

I . I ' I J M

l . I I

. . . . . . . TO+TIIL . I l l . . : ? k * " l i e

(B) M ~ = 4 4 , I.E. C02

, : ; . - - . - ,,, _ . . . . . - _ _ • ,.+ . . . . . _ ..... + +

III III :HI 411 SID 41,11 "PU [ ,q, e+ I . IT~tRT r162 :__ IL_I . ~ CYC:LI: :. IP~!O .. .

ID.3( IBm I N T . R A ~ , [ 1 [ - ~ I

�9 . , , u (C) M ~ = 18, I.E. H 2 0

�9 .,-/ \, /

l �9 ~0 3 e Lira ~ 0 4 0 ?ID C~ ln~ l + + TOTML t t ~ : "+. ? + ~ O , _ S l . . t . . . . . . ST_mm+T C + C L [ :_ I . . . . ~ r .:_ e 4 l . . . . .

Q . & E l l U INrT. Re~:~E- r l l : - I ~ 1

r e , s l i m

�9 ,,,,,, (D) M/Z=28, I.E. CO

B. 4 B I O

I . , q I Q O

I n I Q + 2 ~ ~ ~ 6 1 ~O [ ~ I n J

_ _ T O t l k T ! ! . " ? 8 , M , , O n . 1 2 _ I . . . . . I T M t r162 : +41 . . . . H CY.CU~" | . !

Figure 10.21 MS intensityltime curves for four different M/Z values as measured during the TGA experment with the maize-based polymer sample

374

Figure 10.22 shows an average FTIR gas-phase spectrum measured at TGA temperatures between 200~ and 202.5~ after subtraction of the dominant C02 and H20 absorption effects. The strong absorption maxima at 932 and 966 cm-I confirm the release of NH3 as detected already with the MS.

Also the strong absorption maxima at 2253 and 2285 cm-i are pointing at the presence of R-N=C=0 isocyanate groups in the TGA purge gas. These components were not detected by the MS (only a weak m/z = 43 maximum i.e. possibly H-N=C=O), possibly due to fragmentation of these components in the MS.

The release of components like NH3, CO2 and R-N=C-O due to thermal degradation processes between about 150"C and 250~ from the maize-based polymer sample were not measured for the starch reference sample. This indicates that this maize-based polymer is in fact a chemically-modified-maize based polymer.

The case study reports 10.2 and 10.4 have been published before in the Thermische Analyse Bulletin (TAB). TAB is issued four to five times a year by the Thermische Analyse Werkgroep Nederland (TAWN), the Dutch Thermal Analysis Association. The case study reports 10.3, 10.5 and 10.7 will also be published in TAB.

375

O. oos I '

A

| . oo41

O . m l

O , m

I ! I 1 I " w I I

I l l s I ra- I ) 111 slim- I ) M"NmC'O V|blPIt |OInl l

~ISll eel-. I)

l o l l o n - l l l i l t l e t )

l I I l l ~J I

I l l l oam-I

o. OOlO

4ooo iSOO =0oo zsoo a W ~TSo ~50o ,zso ~ooo 75o C m -1

) H v i i i = i t , lone I | | O S e n - l " O.OOil - I ~q l om" - I ) i 9111 e= - - I )

I ZZll i l ore- I

I I ' l l o

I O. OOlO

Figure 10.22 The average FTIR gas phase spectrum (CO2/H20 subtracted) measured between 200~ and 202.5~ during the TGA experiment with the maize-based polymer film sample

- O. 00)!

O, 0010

-O. 0000 4110

376

References

I. Nielsen, Mechanical Properties of Polymers and Composites, New York, 1974.

2. G.V. Vinogradov e.a., J. Pol. Sc., A-2, 1971, 1153-1171. 3. A. von Hippel, Dielectric Materials and Applications, J.

Wiley, New York, 1954. 4. N. Hill, Dielectric Properties and Molecular Behaviour, Van

Nostrand, London, 1969. 5. V. Lanza and D. Hermann, J. Polym. Sci., 28, (1958) p. 622. 6. F. Wurstlin, Ko11. Zeit., ii/, (1967) p. 79.

Index

377

AI

abrasion resistance, 17 absorption, 81, 111, 151

" level, 113 abundance, 204 accuracy, 145, 356 ac properties, 123, 133, 175 acoustic elastometry, 109

" impedance, iii activation energy, 129, 140, 154, 183, 268 additive, 26, 209, 232 ageing, 105, 312 alifatic acid, 47

" polyketone, 77 " amine, 172

alpha- crystallinity, 302 " -olefins, 235

alumel, 61 amine curing agent, 40 ammonia, 209 amorphous, 232, 233, 278, 283

" phase, Ii, 337 " " ageing effect, 314 " " transitions, 95, 312

amplitude, iii anaerobically aged, 105 angle of incidence, Ii0 iii anhydrite curing agent, 40 anisotropy, 79, 325 annealing, 28, 303, 342 annulus, 165 anti-static compound, 172

" " additive, 209, 354 " " epoxy GFR pipe, 177

aromatic oil, 22 " acid, 51 " amine, 172

Arrhenius, 129 " plot, 137, 154, 331

ASB (3-azidosulfonylbenzoic acid), 366 atactic-PP, 26

" vinyl BR, 282 atomic polarisation, 126, 181 attenuation coefficient, 111 axial direction, 179

Bo

backbone, 264 base-line, 16, 97 bending vibration, 209 benzoic acid, 201, 204 beta- crystallinity, 302 bis (2-ethoxyethyl) ether (BEE), 340 blend, 17, 95, 99

378

blowing agent, 105 boiling point, 204 Bolzmann' s constant, 183 bound ammonia, 219

" water, 214 branching, 232 brass disk electrodes, 184 breakdown phenomena, 123 brittle, 14 brownian motion, 126 Buccsi's theory, 181, 192 butyllithium, 17, 282 bulk modulus, 94, 119

CO

cable compound, 163 calibration DSC, 10

" TGA, 61 " TMA, 77 " dielectric cell, 147 " FTIR, 201 " MS, 201

capacitor, 171 capacity, 85, 124 capacitance bridge, 85, 356 capacitive coupling, 175 capillary, 85, 151

" tip temperature, 200 car-tyre, 17, 140, 364 carbon black, 26, 91, 140, 171, 177, 364

" monoxide, 77, 297 Carilon polymer, 77, 297 carbonyl dipoles, 302, 311, 336 carboxylated system, 366 case studies, 3, 339, 374 casting, 152, 155, 173, 222, 276, 339 chain flexibility, 232

" stiffness, 232 char, 268 charge carrier, 124, 173, 175, 183

" current, 133, 177 " process, 171

chemical structure, 230 cis-, 17 cobalt, 18, 282

" phthalocyanine, 209 cohesive force/energy, 232 compatibility, 19 compressional modulus, 94 compression moulding, 309, 312 complex quantity, 128 composite system, 151 condensation effects, 200 conduction, 171 conductive properties, 123

379

conductance loss, 130 connectivity indices, 232 coupled techniques, 188 correlation coefficient, 235, 241, 254, 355 corrosion protection, 158 Coulomb, 124 counter, 113 creep, 230 critical angle, 109 crossbeam ion source, 200 crosslink density, 105, 245 crosslinking, 14, 95, 232, 245 cruse oil, 158 crystallinity, 264, 283 crystallisation temperature (determination), 14, 345

" half-time value, 85 " process 345

crystalline phase (transition), 312, 337 crystallite perfection, 303, 306

" imperfection, 14 " size, 306

crystallinity, 318 cup-shaped electrode, 189 cure exotherm, 40

" reaction conversion, 43 " schedule, 248, 339

curing agent, 40, 177 current density, 124 cyanogen, 214 cycloalifatic amine, 172 cyclohexane, 10, 19

m O

dc properties, 123, 133, 172, 177 " conduction, 147, 334

Debey' s equation, 129 decay, 158 decomposition process, 62, 209

" products, 195 " temperature, 268

deconvolution procedure, 97 degradation product, 222

" onset temperature, 225, 228 density, 105, 147, 232, 302, 309, 310, 354, 356 derivative curve, 204 densification process, 314 DETA (dielectric thermal analysis), 3 4-4'diamino diphenyl methane (DDM), 248, 339 4-4'diamino diphenyl sulphone (DDS), 250 4-4'diamino diphenyl propane (DPP), 250 4-4'diaminophenyl (DP), 250 dicyanediamide, 222 dielectric analysis, 3 dielectric constant, 124, 133, 145, 147, 152, 158, 172, 309, 331

380

dielectric constant, 356 " properties, 123, 131, 327 " strength, 230

diffuse reflectance, 200, 219 diglycidylether bisphenol A (DGEBA), 222, 247, 339 dilatometry, volume-, length-, 77 diphenylolmethane (DPM), 345 dipole, 126, 183

" loss, 130 dipolar relaxation, 336 dipolarisation, 181 discharge current, 133, 163, 189

" process, 171 dissipation factor, 128 dissociation process, 183 distorsion, 113 distribution, 128 DMA (dynamic mechanical analysis), 3, 94, 314, 350 " /DETA, 188

double bonds, 282 drybox, 69 DSC (differential scanning calorimeter), I0, 318, 342, 364 " high pressure-, photo-, modulated-, 3

DTGA curve, 364 dynamic stiffness, 342, 350

E Q

emulsion SBR, 17 elastic response, 94

" (Youngs) modulus, 94, 105, 119 " constants, ii0, 119

electrical grade PVC, 163 " properties of, 327

electric field, 123, 179, 183, 192 electrification time, 166, 170, 354

" voltage, 167, 170, 354 electrode, 85, 133 electronegativity, 126 electronic conduction, 123

" polarisation, 126, 181 electrometer, 127, 131, 189 electrostatic charge, 158, 171

" properties, 177 " safety criteria, 171

embedded wire electrodes, 359 endblock, 184 endothermic effect, I0, 13

" fusion, 29, 345 end-use properties, 230 engineering polymers, 77 enthalphy relaxation, 11 epoxy coal-tar system, 163

" moulding powder system, 359 " powder coating, 40

epoxy resin, 43, 95, 105, 133, 171, 222, 247, 276

381

epoxy coating, 151 " molar mass value (EMM), 247

equilibrium water saturation, 83, 152, 274, 276, 278, 314 ethylene, 77

" glycol, iii evaporation, 72 evolved gas analysis, 195 exothermic process, i0, 222 expansion coefficient, 175, 322 extinction coefficient, 264 extractables, 364 extrinsic space charge polarisation, 181

F O

failure, 230 ferro-electricity, 123 ferromagnetic (Curie) transition, 61 field strength, 124, 177, 179 filler, 231, 233 fixed frequency system, 94 flexural modulus, 312 fragments spectrum, 204 free volume relaxation, 314 frequency, 94, 97, 105, 109, 113, 140, 192, 331, 356 friction, 230 frozen-in stresses, 79, 322 FTIR (fourier transform infra red),3 function generator, 113 functionality, 247 furan ring, 278 fused silica capillary, 196 fusion, 26, 297

GO

gamma relaxation, 312, 317, 331 guarded electrode, 131, 133, 145, 159 gas cell (FTIR), 196 gas constant, 129 gas-phase spectrum, 204, 209 glass-rubber transition, 189, 192, 327

" " " temperature determination, Ii " " " region, 113, 137, 232

glass fibres, 26, 81, 322 glassy state, 175, 232 glycerol, 105, 201, 204 glycidylether of phenol novolack (GEPN), 222 gold electrodes, 327 green polymer, 371 group contribution additivity, 232

HI

Hamon equat ion, 130 Hartshorn equation, 155

382

Hct-value, 303, 309, 311 heat flow shift, 13 heat-flux DSC, 10 heating rate, 14, 62, 69, 97, 131, 183, 185 heat of fusion, ii, 14,

" " reaction, 40, 47 " " vaporisation, 52

helium, 196 heterogeneous catalyst, 26 hexafluoro-isopropanol (HFIPA), 297 hexagonal crystal lattice, 26 hexahydro-phthalic anhydide (HHPA), 251 Hf-value, 14, 16, 29, 36, 231, 264, 277, 297, 302, 309, 345 Hf (i) -value, 342 Hf (max.)-value, 264, 299 high-potential electrode, 88, 159, 163, 184, 189, 354 high voltage source, 131 " " switching unit, 131

hydrogencyanide, 209, 222 hydrogen bonding, 222

" stretch, 209 hysteresis, 115

I I

immersion (water), 83, 109, 159 " time, 159, 167, 170 " transducer, 113

impact improver, 95 " resistance, 97

inclusion, 175 Indium, ii, 77 induced crystallinity (pressure-, shear-, thermally-, polymerisation liquid-), 303 inert atmosphere, 62, 268 injection moulded, 327 inflection point, 89 in-phase component, 95 insulation resistance, 171 insulator, 124 interface, 175 interracial region, 158 internal mixer blended, 19 interferometer, 200 intramolecular, 278 intrinsic space charge polarisation, 181 ion-free water, 274, 314 ionic impurities, 154 ionol, 69 IR heaters, 201 IR spectrum, 204 IR absorptions, 214 IR vibrations, 219 irganox, 69 iron, 61 isocyanate, 105, 145

383

isotactic-PP, 26 " vinyl BR, 282

isothermal crystallisation, 91 isotropic, 79, 342

Ko

kaolin, 81 KBr windows, 196 kerosine, 159 key-parameter, 151 " -properties, 231, 277

Ki-value, 163

L O

laminate, 152, 155 LCP, 342 �9 LCR meter, 131 1.e.c. (linear expansion coefficient), 79, 322 length dilatometry, 77 level detector, 113 life-time predictions, 107 ligand, 209, 214 lithium, 18 liquid displacement method, 356 liquid nitrogen, 111, 131 local mode relaxation, 331 long glassfibre reinforced, 322 longitudinal waves, 109 low-potential electrode, 88, 159, 163, 184, 189

" " " (spring-loaded), 354 loss factor, 331, 334 loss modulus, 94, 105, 312 loss maximum, 154 LVDT (linear variable displacement transducer), 2, 189

MO

mainchain, 233, 282 " " symmetry, 253 " " flexibility, 253

Maxwell-Wagner absorption, 152 mass increase (TGA), 288

" spectrum, 209 " transfer, 230

masterbatch, 47 matrix, 151, 175, 222, 342 MDI-index, 105, 145 measuring capillary, 88 mechanical waves, 109

" properties, 115 melting temperature determination, 14 melt strenght, 350 mercury, 85, 159 mesophase, 342

384

metallocene catalyst, 26 methyl group, 311 methylisobutylketone (MIBK), 339 MgCI2 catalyst carrier, 69 mica, 81, 350 micrometer, 113, 119 mixture of, 111 moisture, 69, 81, 151, 155, 231, 274, 312, 327 molar mass, 232, 253, 299

" thermal decomposition function, 269 " water content, 274, 276

molecular modeling technique, 232 " structure, 232, 247, 253 " weight, 231, 233 " " distribution, 231 " " increase, 293

monoclinic crystal lattice, 26 MS (mass spectrometry), 3 MS spectrum, 204 m/z-value, 204

N I

naphtenic oil, 22 natural rubber (NR), 19 network, 105, 175, 177, 247

" functionality, 245 " chain density, 245

n-dodecane, 54 nickel, 18, 61, 282 NMR, 264 non-linear behaviour, 147 non-polar, 183, 184, 282 n-tetradecane, 201 nucleating agents, 26

" efficiency, 35 nylon 6, 275 nylon 6.6, 81, 239, 256, 266

O I

oil-extention, 19 oligomers fraction, 269

" in PP, 62 oligomer sample series, 299 orientation polarisation, 125, 181, 184, 192 orthorhombic unit cell, 302 oscilloscope, 113 out-phase component, 95 oxidative atmosphere, 268

P I

Peltier elements, 89 perkalloy, 61 permittivity, 128

385

phase transition, 230 phenolic-OH curing agent, 40

" " residue, 222 phenoxy resin, 240 photo voltaic properties, 123 piezo-electricity, 123 pKa-value, 47 plasticiser, 95, 151,. 167, 231, 233, 314 platinum, IIi, 159 Poisson ratio, 94, 119 polarisation, 123, 181, 183, 189 polarisability, 124 polarity, 232, 253 polar vapour, 204 poly-acrylonitrile (PAN), 240, 256 poly-alkylstyrene, 235 poly-arylene sulphone, 241 poly-butadiene (BR), 17, 115, 256, 262, 266, 271, 282, 294 poly-l-butene (PIB), 38, 235, 238, 256, 262, 266, 271, 292 poly-butene-terephthalate (PBT), 256, 262, 271 poly-carbonate (PC), 240, 256, 262, 271, 275

. " dichloro- , 240 " " , tetramethyl- , 240 " " tetrachloro- 240 " " , tetrabromo- , 241

poly- (cis-chloroprene), 238, 256 poly- chloro-paraxylylene, 239 poly-ether ether ketone (PEEK), 237, 240, 256, 262, 266 poly-ethylene (PE), 238, 261, 266, 271, 275, 294, 356 poly-ethyleneoxide (PEO), 261 poly-ethylene-propylene rubber (C2C3 rubber), 97, 238 poly-ethylene-terephthalate (PET), 239, 256, 262, 266, 271 polyethylene- succinate, 239 poly-formal, 240, 256, 266 polyformaldehyde, 256 poly-l-heptene (PIH), 237, 238 poly-l-hexene (PIHex), 238 poly-isobutylene (PIB), 238, 256, 262, 271 poly-isobutylene oxide, 238 poly-isoprene (IR), 85, 238, 256, 271, 282, 288, 294 polymerisation process, 184 poly-methylacrylate (PMA), 235, 239 poly-methylmethacrylate (PMMA), 240, 256 poly-3-methyl butene-1, 239, 261, 266, 271, 294 poly-4-methyl pentene-l, 239, 261, 266, 271 polyketone (PK), 77, 297 polyol, 105, 145 poly(l-olefin) s, 36 poly- oxyisobutylene, 256 poly-oxymethylene (POM), 238, 256, 266, 271 poly- oxyt rimethylene, 238 poly-oxy-paraphenylene (PPO), 240, 256, 262

" " " , dimethyl- , 240, 261 poly-paraxylylene, 239 poly-l-pentene (PIP), 38, 238, 256, 262, 292 poly-pivalolactone (PPL), 256, 266, 271, 275

386

poly-propylene (PP), 14, 26, 38, 64, 235, 239, 261, 266, 271, 275, 292, 342, 350

poly-propyleneoxide (PPO), 261 poly-styrene, 240, 256, 266, 271, 275, 294

" " foam, 354 poly-styrenebutadiene (SBR), 17, 115, 140, 184, 245 poly-sulfone, 241, 275 poly-tetrafluorethylene (PTFE), 196, 261 poly-trans 2.3 epoxybutane, 266 poly-thio-paraphenylene, 240, 256, 262 poly-vinylacetate (PVAc), 239 poly-vinylalcohol (PVA), 239, 256 poly-vinylchloride (PVC), 151, 163, 175, 184, 192, 239, 256, 275 poly-vinylfluoride (PVF), 239, 256 poly-vinylidene fluoride (PVDF), 238, 256 poly-vinylidene chloride (PVDC), 239, 256 poly-vinylpropionic acid, 239 poly-urethane (PU), 105 postcured, 133 power-compensating DSC, 10 precision, 13 precursor, 209, 214 processing procedure, 308

" window, 231 propagation direction, 109

" speed, Ii0 propylene, 77 pulse, 109, 113

" generator, 113 " length, 113 " travelling time, 113

purge gas DSC, 10, 54 " " TGA, 61 " " TMA, 77 " " TGA, 196

pyridine, 225 pyro-electricity, 123

O

quadrupole analyser, 200

RO

radial direction, 177, 179 reaction exotherm, 47 recrystallisation temperature, 14, 26 receiver, 109, 119 recombination process, 183 reflection, 109 relative humidity, 147 relaxation effect, 95, 123, 141, 331

" time, 129, 158, 171, 183 repeatability, 13, 16, 167, 325 reproducibility, 13, 16, 89, 318 residual monomer, 231, 233, 253

387

residual reaction exotherm, 47 " solvent, 233, 253, 340

resistance thermometer, 131 resistive behaviour, 175 resistivity (volume), 3 resonant system, 94 resonance frequency, 113 rigid, 181

" PU foam, 105, 145 " amorphous phase [x(r,a)], 321

rolling resistance, 17, 113, 140 round robin test, 13 rubber, 17, IIi, 113, 181, 245, 282

" compound, 91 rubbery state, 175, 232, 282

" plateau, 350

So sample apertures, 111

" dimensions, 77 " disks, 111, 133, 145, 189, 327 " high pressure pans, 47, 52 " holder, 87, 111, 189 " loading procedure, 69 " " table, 72 " nib, 64 " pans, 14 " pill, 105 " sheet, 322, 325 " strips, 274 " weight, 14, 204

screen, 88, 189 seawater, 158 secondary relaxation effects, 181, 312 self-reinforcing composite system, 342 self-seeding effect, 16 semi-crystalline, 14 " -static Td(o)-value, 69, 269

shear modulus, 94, 119 " waves, 109, ii0

shift factor, 141 short-circuited, 127 shortcircuiting, 171, 179, 184 shrinkage, 77, 81, 322 side-chain crystallisation, 36, 235 side-group, 235, 253, 282 signal distorsion, 113 silicone oil, 111 silver effect, 189

" electrodes, 134, 189 simultaneous techniques, 3, 189 sinusoidal, 128 solution blended, 19 solvent in a thermoharding system, 339 SIS sequential block copolymers, 290 space charge polarisation, 183, 192

388

specific absorption, 209 " heat, 95, 264, 318, 321 " material properties, 337 " thermal expansivity, 17 " volume, 91, 264 " volume resistivity, 124, 131, 133, 158, 163, 172, 177,

327, 334, 355 " surface resistivity, 355

spherical interstices, 155 spherulite size, 26 spherulitic growth rate, 264 sprue-plane, 79 stability, 189

" DSC, i0 stainless steel capillary, 196 steric effects, 232 step-wise heat flow shift, 13 stoichiometric, 43, 248 strain, 94 stray capacitance, 147 stretch vibrations, 214, 219 stress, 94 surface resistivity, 354 surface tension, 177 susceptibility, 129 symmetrical polymers, 253 syndiotactic- PP, 26

" vinyl BR, 282, 286 synthesis, 209

TD

tacticity, 232, 253 tail structures, 184 talc, 26 tan delta, 95, 105, 128, 133, 140, 147, 152, 309, 331, 342 tankcoating system, 158 tapered SSBR, 184 tapwater, 163 TA (thermal analysis), 2 Tc-value, 14, 16, 28, 36 Tc (i) -value, 342 Td(0.5)-value, 62, 268 Td (o) -value, 62, 268, 278 Te-value, 14 TGA (thermogravimetric analysis), 3, 268, 288, 364 " /DSC, /DTA, /FTIR, /GC, /MS, 188

TGA - coupled - FTIR/MS, 195, 369, 371 Tg-onset value, 13, 172, 231, 277, 283, 286, 290, 314, 318, 340,

"342 Tg-value contribution per crosslink, 250 " " determination, ii, 318 " " increase, 293 " " of, 17, 36 " " " crosslinked epoxy resin systems, 249 " " " crosslinked polymeric systems, 245

389

Tg-value of vinyl polymers with linear sidechains, 235 Tgl-, Tg2-value, 51 Tg (midpoint) -value, 77 Tg/Tm correlation, 254 thermal expansion coefficient, 77, 189

" history, 16, 28, 253 " insulation, 105 " stability, 61 ,231, 268 " transitions, 181

thermally stimulated current, 181 thermo-luminosence, -magnetometry, -optometry, - sonimetry, 3

" -electrometry, 123 " -dilatometry, 77

thermogram, 13, 30, 31, 34, 37, 40, 41, 42, 48, 55, 57. thermosetting resin, 40 thermostat bath, 111, 131 Ti-value, 342 TIC14 catalyst, 69

,' .DIBP complex, 69 time, 94, 141

" /temperature superposition, 140 " dependency, 177

tin, 11, 18 Tin-, Tml-, Tin2- (value), 14, 16, 28, 36, 231, 253, 277, 286, 297,

302, 309, 345 Tin-value depression, 297 Tm(o)-value, 254, 299 Tm'-value, 303, 309, 311, 342 TMA (length -, volume - thermomechanical analyser), 2, 3, 322 " /DETA, 188

T-onset, II torsion pendulum apparatus, 94 toughened PP, 95, 97 toughness, 230 tracking, 230 trans-, 17 transducer, ii0, 113 transfer line, 196 transmitter, 109 transmission, 109 transition, 175 transverse waves, 109 tribo-electricity, 123 triclinic crystal lattice, 26 trigger signal, 113

" unit, 113 trimethylpyridine, 225 TSDA (thermostimulated discharge analysis), 3, 123, 181 TSD/TMA, 189

UO

Udel, 240 UL-index, 268 ultimate properties, 230 ultra-high molecular weight PP, 350

390

ultra-low heating rate TGA, 225, 269 ultra-sonic beam, 113

" " loss factor, iii " " measurement s, 109 " " properties, Ii0

Underwriters Laboratories Testing Procedure, 268 unsymmetrical polymers, 253

VQ

vertical furnace TGA, 196 vibrate, 109 vibration, 204 vinyl-, 17, 235 viscous response, 94 voltage, 166

" supply, 189 volume dilatometry, 85 volume resistivity, 354 vulcanizate, 85, 245 vulcanised rubber, 364

WO

water absorption, 151, 274 " clusters, 158

wavelength, 109, 204 wear, 230 weight loss, 268 wet deposition technique, 322 wet grip, 17, 113, 140 WLF-equation, 141 wollastonite, 81

X. x (c) -value, 300, 302, 309, 310, 321, 350 X-direction, 322, 325 x-ray diffraction (XRD), 297 x-ray scattering (XRS), 264 XRD spectrum, 303

Y" Y-direction, 322, 325 yield strength, 230

Z Q

Z-direction, 322, 325 zero, 126, 147

Documents

Characterisation Of Polymers By Thermal Analysis - W Groenewoud (Elsevier, 2001) Ww