A study of the ultra-violet degradation embrittlement of polypropylene polymer

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A Study of the Ultra-Violet Degradation Embrittlement

of Polypropylene Polymer

Proefechrift ter verkrijging van de graad van doctor

aan de Technische Universiteit Delft,

op gezag van de Rector Magnificus,

prof. drs. P.A. Schenck,

in het openbaar te verdedigen

ten overstaan van een commissie door het

College van Dekanen daartoe aangesteld,

op dinsdag 25 oktober 1988, te 16.00 uur

door

G.E. Schoolenberg

geboren te Rotterdam,

ingenieur Industrieel Ontwerpen.

Dit proefschrift is goedgekeurd door de promotoren:

Prof. ir. J.L. Spoormaker Prof. dr. ir. A.K. van der Vegt

Colofon:

ISBN: 90-9002498-0 .

Delft University of Technology

Department of Industrial Design Engineering Section Engeneering Design

Technische Universiteit Delft

Faculteit van het Industrieel Ontwerpen Vakgroep Konstruktie

Druk: Drukkollektief Luna Negra bv, 1988 Omslagfoto: A.M.W.J. Bakker

Preface and Acknowledgements

In the research described in this thesis several disciplines (fracture mechanics, materials science, polymer physics and organic chemistry) meet. A variety of experimental techniques and theories from these disciplines was applied. In such a research it is indispensible to consult with specialists, who are prepared to make their knowledge easily accessible to and applaud its application by a colleague from another field.

I was happy to co-operate with institutions, where such specialists could be consulted, which is proven by the long and even very incomplete list of acknowledgements presented below. This thesis could not have been written but for their ready support and that of many other people, to whom I would like to express my gratitude.

First of all I will mention my collegues from the Laboratory of Mechanical Reliability in the Faculty of Industrial Design Engineering especially Mr.H.D.F.Meyer for carefully and reliably performing the experiments and Jhr.Ing.D.C.H.Smissaert, for his creative computer programming. Furthermore, both made countless useful contributions to the experimental set ups and procedures.

Mr.M.W.van Dalen, who proved to be a very supportive colleague through all the ups and downs of the past years and who also made the nice diagrams, deserves special thanks, as does Dr. D.J.Cumberland for editing the english text.

The research was performed in conjunction with the Plastics and Rubber Research Institute of TNO in the Netherlands, and I am much indebted to its previous manager, Prof.Dr.Ir.L.C.E.Struik. During one months stay I became familiarised with many experimental techniques, which proved to be very useful.

I am especially indebted to Dr.Ir.D.J.van Dijk and Dr.P.Vink from this institute for many valuable discussions and suggestions. Also Ultra-Violet degradation facilities were kindly made available and the specimens were treated with great care by Mr.W.Boels.

The polymer material was supplied by Dutch State Mines, who also advised with respect to mechanical test facilities. Many other services were rendered by Mr.W.Leunissen of DSM, for which my thanks.

Microtoming techniques were taught to me by Ing.P.Wildervanck of the Plastics and Metalware Factory of Philips in Eindhoven. To Prof.Ir.A.Anemaat of this company I am indebted for the moulding of the specimens.

The Materials Science and Chemistry Faculties of the University of Delft were very helpful. I am especially obliged to Prof.Dr.Ir.A.K.van der Vegt and the members of his research group for their hospitality and assistance. Also I am grateful to Ing.E.J.A. van Dam who made the many scanning electron micrographs, and to many others in the Metallurgical Laboratory.

Contents

page

List of symbols used 13 Introduction

Aim of the research 15 Lifetime Prediction of Plastic Products 15

under Outdoor Conditions : Present Situation

Failure Mechanisms: Yield and Fracture 17

1 Literature Survey

1.1 The Process of UV Degradation of Polymers 19

1.1.1 UV Absorption and Photoinitiation 20 1.1.2 Photooxidation: a Free-radical Chain 22

Mechanism

1.1.3 Photodegradation : Chemical Reaction 23 Processes in PP

1.1.4 Chain Scission and Cross Linking 24

1.1.5 Molecular Mass Reduction 26 1.1.6 Effect of other Environmental Conditions 27

1.1.7 Stabilisation 27 Conclusions 29

1.2 Loss of Mechanical Properties

Introduction 30 1.2.1 Thin Films : Correlation beteen Chemical 30

and Mechanical Test Results

1.2.2 Thick Specimens : Surface Defects 33 1.2.3 Thick Specimens : Tensile Tests 35 1.2.A Thick Specimens : Impact Tests 35 1.2.5 Thick specimens : Flexural Properties 36

Conclusions 36

1.3 The Behaviour of Bodies with Surface Cracks

Introduction 37 1.3.1 Linear Elastic Fracture Mechanics-.Theory 37

1.3.2 Failure Stress - Crack Length Relation 38 1.3.3 Tensile Behaviour of Degraded Specimens 39 1.3.4 Impact Behaviour of Degraded Specimens 41

Conclusions 42

References 43

2 Scope of the Current Research

Introduction ^7

2.1 Test Procedure for Checking the Fracture 48 Mechanics Approach

2.2 Chosen Test Parameters 49

2.2.1 Processing Conditions and Stabilisation 49

2.2.2 Specimen Geometry 50 2.2.3 Temperature and Deformation rate 50

References 50

3 Mechanical Test Facilities and Methods, Test Materials

Introduction 51

3.1 Mechanical Test Methods

3.1.1 Three Point Bending Test 51

3.1.2 Tensile Impact 54 3.1.3 Notching 54 3.1A Data Processing 55

3.2 Specimen Preparation, Morphology and Initial Properties

Introduction 56 3.2.1 Compression Moulding 56

3.2.2 Injection Moulding 57

3.2.3 Stabilisation 57 3.2A Initial Properties 58

Conclusions 64 References 64

4 Failure Properties of Degraded Specimens

Introduction 65

4.1 Artificial Weathering in the Xenotest 1200 65

4.2 Failure Properties in Impact 66

4.3 Failure Properties at v-0.01 m/s 69

4.4 Comparison of the Different Specimens 72

4.5 Failure Properties at Smaller Specimen Depth 73

Conclusions 75

References 75

5 Depth of Embrittlement

Introduction 76

5.1 Visual Inspection an Microtoming 76

5.1.1 Appearance of the Degraded Surfaces 77

5.1.2 Cross-sections 78 5.1.3 Microtome Slices Parallel to the 80

5.2 Infra-Red Spectrometry

Introduction 82 5.2.1 Instruments and Specimen Preparation 82

5.2.2 Analyses of the Spectra: Carbonyl

Formation 83 5.2.3 Carbonyl Formation : Results 84

5.2.4 Analyses of the Spectra: Helical 87 Content, Crystallinity

5.3 Relation of Carbonyl Content to Brlttleness

5.3.1 Fracture Surfaces 89 5.3.2 Stress-strain Curves of Microtome 90

Slices

5.4 Summary of the Results on Depth of Embrittlement, 94 Discussion

5.5 Review of Literature Data and Models

5.5.1 Oxygen Diffusion Limitation 97 5.5.2 Comparison to Data on PP 99 5.5.3 Variation of the UV intensity with Depth: 100

the Effect of Stabilisation

5.5.4 Difference between Compression and

Injection Moulded Samples: the Effect of

Morphology 104

Conclusions 105 References 106

6 Failure Properties of Notched Specimens

Introduction 108

6.1 Some General Aspects of Fracture Mechanics 109

6.1.1 Linear Elastic Fracture Mechanics 109 6.1.2 Crack Tip Plasticity and Specimen 109

Thickness

6.1.3 Crack Size and Remaining 111 Ligament Requirements

6.1 A Other than Linear Elastic Material 112

Behaviour

6.2 Linear Elastic Approach, G jc determination 115

6.2.1 Plati-Williams Analysis 115 6.2.2 Preliminary Results: 116

Linear Elastic Approach

6.3 Linear Elastic Approach, K ic determination 118

6.3.1 Preliminary Results: Kjc with and without 118

Plastic Zone Size Correction

6.3.2 Estimation of the Plastic Zone Size 118 and Size Requirments

6.4 Qualitative Description of the non-LEFM Behaviour

Introduction 122

6.4.2 Plastic Hinge Formation and

Non-Plane Strain Behaviour 122

6.4.2 The Relationship Between K and G 125

6.4.3 Viscoelastic Effects 127

Causing Stable Crack Growth

6.4.4 Relation to Data on PP in the Current Research 131

6.4.5 Fracture Surfaces 132

6.4.6 Discussion 133

6.4.7 Relation between K and G: 136

Effect of Load Time and Deformation Rate

Conclusions 138

6.5 Review of Literature Data 139

References 141

7 Validity of the Fracture Mechanics Approach

Introduction 143

7.1 Comparison of Notched and Degraded Specimens

in Impact (1.5 m/s) 143

7.1.1 Failure Energy 143 7.1.2 Failure Stress 146

7.2 The Effects of Deformation Rate 146

7.2.1 Specific Failure Energy 146 7.2.2 Failure stress at 0.01 m/s 149 7.2.3 Varying Deformation Rates 150 7.2.4 Classification of the 151

Degradation Behaviour

7.3 Evaluation : Fracture Processes 153

7.3.1 Relation between Fracture Path, Fracture

Surface and Specific Failure Energy 153

7.3.2 Fracture Processes 155 7.3.3 Fracture Energy-Degradation Time 160

Relation : Conclusions

7.3.4 Conditions for the Crack Speed Effect 160 Conclusions , 163

7.4 Further Aims of the Investigation 164

References 165

8 Tests below th Glass Transition Temperature

Introduction 166

8.1 Fracture Toughness

8.1.1 Linear Elastic Approach 166 8.1.2 Size Requirements and Plastic Hinge Formation 167

8.IA CTOD, Mirror Zone, Fracture Surfaces 168

8.1.5 Crack Length Dependence 168

Conclusions 171

.2 Comparison of Degraded and Notched Specimens

Introduction 172 8.2.1 Failure Properties of Degraded 172

Specimens at -A5°C

8.2.2 Comparison of Degraded and 174 Notched Specimens at -45°C Conclusions 174 Reference 175 9 Pre-cracking by Fatigue Introduction 176 9.1 Fatiguing Procedure 176 9.2 Results 177 Conclusions 178 10 Tensile Tests Introduction 179

10.1 Fracture Toughness Parameters in Tension 179 10.2 Comparison of Notched and Degraded Specimens 179

Conclusions 181

Summary and Conclusions 183 Nederlandse Samenvatting en Conclusies 185

Appendix 1 Dimensions of the Test Specimens

Appendix 2 Constructional and Experimental Details of the Three Point Bending Tests

Appendix 3 Other than Fracture Related Events in the Three Point Bending Test

Appendix 4 Failure Criteria for the Three Point Bending Test Appendix 5 Tensile Impact Test

Appendix 6 Fractography of Notched Specimens Appendix 7 Fractography of Degraded Specimens

L i s t of symbols used

6 = error in control system E = strain E f = strain at failure Er — strain at b r e a k or rupture Ey = strain at y i e l d p = materials density 0 - stress Cc - critical stress

Of = failure stress, flexural strength

Gr = ultimate strength

Gy = yield strength

t = integration time constant (in hydraulics)

t0 i = characteristic time (visco-elastic behaviour)

<p = angle o f rotation

<p = geometrical factor accounting for finite size in fracture mechanics U - frequency of light u = angular frequency mV,mm % % % % g / c m3 N / m m2 N / m m2 N / m m2 N / m m2 N / m m2 m s m s 8-1 rad/s a = crack length, coefficient o f absorption a0 = defect size, inherent flaw size

a^ - degradation depth

5 «■ average number o f scissions p e r chain A a - length o f mirror zone

è — crack velocity B - width, thickness

c = plastic constraint factor c' = materials density

c0 - oil column stiffness

c^ - hydraulic stiffness cs = specimen stiffness

C = specimen compliance, oxygen concentration C ( t ) = compliance function

time dependent specimen compliance D = depth,

- oxygen diffusion coefficient E = Young's modulus

G = energy release rate

Gjc - c r i t i c a l energy release rate, tensile opening mode

h = Planck's constant I = light transmitted Ia - light absorbed

I0 - incident light

Iy = second moment o f area L = length k = rate c o e f f i c i e n t - c o n s t a n t c m2/ g |im, mm |im, m m m m m/s m m g / c m3 N/m N/m N/m m / N m l / g m2/ N m / N m m cm'/s N / m m2 k J / m2 k J / m2 cd cd cd m m ^ m m

K «Ie KC2

t?

H

critical stress intensity in tensile opening mode

plane stress critical stress intensity candidate plane strain fracture toughness length

power law index

slope of the triangular input signal, tup-velocity

bending moment

number average molecular mass viscosity average molecular mass weight average molecular mass mass

power law index power law index oil pressure load

fully plastic limit load power law index

flow rate input signal

fracture resistance, crack tip radius length parameter

size of the plastic zone correction (Irwin) Dugdale plastic zone size

time

time for crack initiation oscillation period

ductile brittle transition temperature (ace. ASTM D746)

glass transition temperature melting temperature

stored elastic energy

surplus energy released during crack extension per unit thickness kinetic energy dissipated during crack extension per unit thickness

potential energy

total amount of energy dissipated in fracture displacement

crack tip opening displacement velocity

coefficient of viscous friction fracture energy

specific fracture energy (per fractured area) displacement, distance

geometrical factor accounting for finite size in fracture mechanics

load applying at control system

MPaJ"m MPaJ"m MPaJm mm mV/s m/s N/m k g b a r N N dm-vmin mV k J / m2 |im \im mm mm h , s , ms h , s s °C °C, K °C, K k J , J J/m J/m k J , J k J , J m, mm lim.mm m/s kg/s kJ.J.Nm kJ/m2

Introduction

Aim of the Research

This research is concerned with the effects of weathering on the mechanical reliability of plastic products. Under outdoor conditions, polymers degrade, mainly due to UV-light, while heat and humidity have secondary effects. Degradation reduces the mechanical reliability, and in some cases this reduction is very drastici-'-] .

The ultimate aim of this research is to develop a method, which uses mechanical concepts for the prediction of the reliability of plastic products, which are used under outdoor conditions. Such a method should allow a well-founded choice of design parameters, such as type of polymer, dimensions and geometry, which guarantees load bearing abilities over the projected life-time.

The current methods for predicting the loadability of plastic products subject to degradation are discussed in this introduction, and

in the literature survey (Chapter 1 ) .

The usual procedure in predicting mechanical reliability, is to draw up a calculation method which yields a quantity for the load applied and for the loadability of the material used. Comparison of both quantities determines whether the product will fail or not. Often the loadability decreases during the lifetime, due to chemical, physical or mechanical processes. This is the case with degrading plastic products. Lifetime predictions then depend upon the knowledge of loadability as a function of time. In this research the loadability as a function of time was expressed in terms of a fracture mechanics concept, which will be treated in Chapter 2.

The chemical degradation process due to UV-light is very complex. The predictability of the loadability depends on the predictability of the chemical process, which is affected by many side-effects. Chemical processes are not the issue of this thesis. The research has concentrated on the effect of a given type of chemical degradation: accelerated UV degradation in an artificial weathering machine. Therefore the validity for outdoor use of the quantitative statements that result from this research, is limited to cases where the chemical process,^though it may be slower, is similar to that occurring in artificial, accelerated conditions. The study of the chemical processes was limited to the aspect that was the most important for the fracture mechanics approach, the penetration depth of the degradation process in solid polymer parts.

Though the research did not give the ultimate prediction method for the loadability of plastic products, it did reveal the relation between chemical effects and the mechanical loadability.

Lifetime Prediction of Plastic Products under Outdoor Conditions: Present Situation

When polymers are weathered, the loadability decreases. Both the short term and the long term mechanical properties are affected. This research is mainly concerned with the short term properties.

In. time the ultimate strength is reduced. This reduction is particularly severe in impact. The failure mechanism changes from yield

to fracture, under the same loading conditions, and in short, the material embrittles. (N.B.This embrittlement must not be confused with that caused by the so-called physical ageing phenomenon, which is less severe.)

The reason for this change in failure mechanism is still obscure from the mechanical point of view. It is not clear which are the relevant material parameters, and in consequence it is often questionable how loadability should be tested, let alone calculated. This renders quantitative reliability predictions impracticable, while statistical reliability analysis is inconceivable at the moment. Qualitatively the method illustrated in fig.l.i is usually applied for lifetime predictions of the polymer m a t e r i a l ^ ] .

A mechanical test, e.g. a tensile or impact test is applied to the degraded specimens. The elongation, stress or energy at rupture is measured as a function of degradation time. Degradation usually takes place in an artificial weathering machine or in extreme climates, where the testing can be accelerated. The lifetime is then estimated as the time it takes for the measured quantity to be reduced to half its original value (the half-time), multiplied by the acceleration factor of the test environment relative to the real environment. It is clear that the relevance of this lifetime to the real lifetime, depends on the shape of the curve in fig.1.1, the level of loading in practical use, the accuracy of the acceleration factor and the appropriateness of the chosen test method. Implicitly it is assumed that the loadability after the half-time is so much reduced, that the half time can be used as a failure criterion.

100%-

50%-F i g . 1 . 1 Half-time c a l c u l a t i o n of mechanical t e s t parameter

Because many polymers have the type of behaviour above ( i . e . stable

i n i t i a l l e v e l , sharp decrease, very low end level) t h i s method i s

suitable as a rough indication of the lifetime.

Predicting more precisely when and at what load fracture will occur however, is only possible after mechanical testing of material of exactly the same composition, that is exposed to exactly the same conditions as are encountered during the service life. This procedure is usually too expensive and time consuming. And even then, the appropriateness of the mechanical test method has to be considered carefully. It is possible that the failure mechanism changes from yield to fracture at another time in the test-circumstances chosen (temperature, deformation rate) than under the conditions of application.

Therefore, the designer usually has to be content with information about the degradation resistance, relative to that of other materials rather than absolute data on the material chosen.

Failure Mechanisms : Yield and Fracture

The degradation process and the way it affects the failure mechanism have to be investigated, before the reliability as a function of time can be described. So far it has been established that in thicker parts only a thin surface layer is actually degraded and eventually causes fracture of the entire part. It is generally assumed that the embrittled surface of degraded polymers, which cracks at low loads, or even spontaneously, has an effect which is similar to the introduction _of_a_cracki^^8-=-iJ—(-see_f ig-.-l-.-2

-)-.—This—implies—that—the—steep-decrease—in-mechanical properties occurs when the surface is sufficiently degraded for surface cracks to be formed which change the fracture mechanism from net section yield, to fracture by crack extension. The validity of this hypothesis will be investigated in this research.

g degraded

Fig.1.2 Analogy between degraded material and material with a surface notch

Yield and fracture are different failure mechanisms for which different calculation methods have been proposed and applied so far. The Tresca or Coulomb criterion, or the von Mises-Huber-Hencky criterion are usually applied to predict the onset of yield. On the other hand the Rankine criterion, or the more sophisticated criteria supplied by fracture mechanics theories are used for the prediction of fracture. Yield criteria suppose a homogeneous, ductile material, failing by plastic deformation, whereas fracture mechanics supposes a brittle material, containing a flaw, and failing by fracture due to crack extension.

Attempts to bridge the gap between yield and fracture based failure criteria were made from either side. These developments are relevant to the degradation problem, because the failure behaviour changes from yield to fracture during degradation.

The effect of discontinuities on local stresses, and their effect on strength on the one hand, is studied and described by stress concentration factors!^] and the material's notch sensitivity respectively.

More relevant to the degradation problem however, are fracture mechanics studies, aimed at predicting crack growth initiation in more ductile materials, involving local yielding. The elasto-plastic fracture mechanics theories show several complexities such as size effects and material dependence. With some adaptions they can be used satisfactorily for establishing fracture properties of ductile materials, such as many

polymers. [->J po r failure prediction however, they cannot be considered

ready for practical use in many cases.[°J

Despite these drawbacks, in the current research the possibilities of applying fracture mechanics criteria to the embrittlement of polymers due to weathering have been explored, because it is the most promising approach. Polypropylene (PP) was chosen as a representative material, because it shows surface degradation and a sudden change from ductile to brittle failure.

References

[1] Dolezel.B., "Die Bestandigkeit von Kunststoff und Gummi", Carl Hanser Verlag, München, 1978

[2] Hawkins,W.L., "Polymer Degradation and Stabilisation", Springer Verlag, Berlin, 1984

[3] Runhnke.G.M., Biritz.L.F., Plast.Polym., 6 (1972) 118-124

[4] Neuber.H., "Kerbspanningslehre", Springer Verlag, Berlin, 1985

[5] Williams,J.G., "Fracture Mechanics of Polymers", Horwood, Chichester, 1984

[6] Broek,D., "Elementary Engineering Fracture Mechanics", Martinus Nijhoff Publishers, Leiden,1986

1. Literature Survey

This chapter summarises literature on the subject of this thesis. The literature survey concerns three topics.

First, relevant physical and chemical information on the process of UV degradation of polymers is given in section 1.1. Special attention will be paid to polyolefins, because polypropylene (PP) was selected for this research.

The second topic covers the effect of the UV degradation process on the mechanical properties of polymers, which is treated in section 1.2. Thin films and thick specimens are treated separately, as increasing thickness causes UV degradation to become more of a surface phenomenon. This research is mainly concerned with thicknesses over 2 mm, because they occur in constructional parts of interest.

The third topic, the mechanical behaviour of bodies with (surface) cracks, is treated in section 1.3. In the same section the inter­ relation between this subject and that of section 1.2 is commented upon.

1.1 The Process of UV Degradation of Polymers Introduction

This section gives information on the physical and chemical aspects of the UV degradation process of polymers. It is limited to those subjects that are relevant for stating the problems of our research, or that account for the test methods used further on. The chemical aspects of UV degradation will be treated according to scheme 1.1, where the steps of the process are listed in logical order. These steps are covered in section 1.1.1 through 1.1.5. Effects of environmental conditions other than light are mentioned in section 1.1.6 and finally stabilisation is treated in section 1.1.7.

subject section UV absorbance (1.1.1) I photolnitiation (1.1.1) I photooxidation (1.1.2;1.1.3) I

chain scission and cross linking (1.1.4) I

reduction of molecular mass (1.1.5)

i

toughness reduction (1.2)

Scheme 1.1 : Steps in the UV degradation process

All organic polymers are to a greater or lesser extent degraded upon exposure to sunlight. The process of degradation by light, either

natural or artificial, is called photodegradation. Degradation under

usually termed weathering. A classification of polymers by their

degradation resistance, without being stabilised (i.e. protected against

degradation by chemical additives), results in three classest^l. Polymethylmethacrylate (PMMA) and polytetrafluorethylene (PTFE) retain their properties for years outdoors, and are used without a stabiliser. Moderately stable polymers such as polycarbonate (PC) and polyethylene terephthalate (PETP) can be used without stabiliser, but often stabilisers are advisable. Other polymers, for example polypropylene (PP), polyethylene (PE), polystyrene (PS), polyamide (PA), cellulose and natural rubber, have poor stability and need extensive stabilisation for outdoor use.

The demand for inexpensive, degradation resistant polymers, stimulated extensive study of thermal and photodegradation processes. In the past forty years a vast quantity of literature has accumulated. The research was aimed at understanding the process, and finding means to prevent it, or to slow it down.

The book of Ranby and Rabekt^J is considered a classical work on

this subject. It contains a review of the knowledge of chemical aspects of photodegradation and stabilisation at that time (1975). More practical information on the weathering behaviour of polymers can be found in the extensive literature survey of Dolezelt^] . Here, not only chemical changes, but also changes in mechanical properties and appearance of non-stabilised polymers and the effect of many stabilisers and pigments are described. In the past few years, developments in weathering, or more specifically photodegradation and photostabilisation

of polymers, were outlined in several publications l^~°\.

The photodegradation and photostabilisation of polyolefins (especially PE and PP) were studied most extensively. These are inexpensive polymers and widely applied in outdoor conditions. However, their intrinsic stability is low, and stabilisation is indispensable. Reviews on the developments in the understanding of their photodegradation process and methods for their stabilisation were published recently!'""J.

Pigments and dyes play an important part in polymer stability. For a long time they have been applied in polymers more or less on a trial and error basisl^J. Though much is known about the photochemistry of pigments and dyes, their interaction with the polymeric environment has not been studied until a few years agot^-lA] .

1.1.1 UV Absorption and Photoinitiation

UV-light is considered the most important component of the weather causing chemical changes. A photochemical reaction is possible when two conditions are met :

- the wavelength of the incident light (i.e. the energy of the lightquanta or photons) is such it can be absorbed by the polymer molecules

- the energy of the absorbed photons is sufficient to dissociate bonds in the polymer molecules.

The first condition is met only if the difference between two energy levels which the molecule can occupy, is equal to the energy of the

photon. Secondly, dissociation of a molecular link occurs when the amount of vibrational energy reaches the energy of dissociation. Bond dissociation by the absorption of light, preceeding the photodegradation

processes, is called photo initiation.

Information on dissociation energies have been published for many organic compounds, amongst others by Calvert and Pitts"-*]. The energy of a photon is E=hv where h is Planck's constant and V is the frequency. The photon has a higher energy when the wavelength is shorter. The higher frequencies are therefore generally the more harmful. Comparing bond dissociation and photon energies, it becomes

clear that UV-light of wavelengths shorter than 400 run [3,6]_ ha s a

photon energy high enough to dissociate bonds in polymer molecules. The relevant wavelengths are between 295 and 400 ran, because wavelengths shorter than 295 nm are filtered out by the atmosphere. The photosensitivity of polymers is due to their absorption of light of wavelengths between 295 and 400 nm.

The photoinitiation of polyolefins poses a problem because pure saturated hydrocarbons do not absorb light in this UV-range. Their photoinitiation was attributed to UV absorbing impurities

(chromophores). The photoinitiation reactions are of primary interest

because they are considered to be the key to a successful development of photostabilisers. The photoinitiating species proposed for PP are ranked in order of importance as follows I°> " 1 :

- PP hydroperoxides (PP00H) and peroxides formed by thermal oxidation during processing

- catalyst residues

- polynuclear aromatic compounds - carbonyl groups

- PP-O2 charge transfer complexes - active oxygen species

- other impurities and additives (pigments and dyes, metal particles, conjugated unsaturation)

A description of all mechanisms by which these photoinitiators give rise to bond dissociation is beyond the scope of this thesis. Furthermore, the photoinitiation process in polyolefins is still under discussion!*.8-9.16-18,22].

After initiation by any of the photoinitiating species mentioned, the subsequent reactions generate hydroperoxides. So even if hydroperoxides are not present in the very beginning of the process they will form quickly, and play a dominant part in photoinitiation after a short time. After longer exposure, carbonylgroups will form in secondary reactions and above a certain amount, they will also become important photoinitiators.

[18]

Some researchers have proposed that carbonyl groupsl2J(especially

in P E I ' Z ] ^ a combination of carbonyl groups as strong chromophores with

hydroperoxides (which are more easily decomposed)

[19]. or PP-O2 charge

transfer complexes I'0] a r e important photoinitiators in the beginning of

1.1.2 Photooxidation: a Free-radical Chain Mechanism

process in Bolland!21] There is general agreement that the photooxidation

polyolefins occurs through a free radical chain mechanism.

proposed the reactions in scheme 1.2 for thermal oxidation, and they are believed to be closely parallel to those in photooxidation.

This reaction process is autocatalytic: the reaction accelerates

itself. After the formation of a free radical by photoinitiation, propagation reactions follow. These are chain reactions which can involve many molecules in the process after only one initiation event.

Initiation üH/hv P -. R- (1.1) Propagation : R- + O2 R O 2 + RH R02-ROOH + R-(1.2) (1.3)

Branching : ROOH -» R0- + -OH 2R00H -» R02-+ R0- + H20 R0- + RH -> ROH + R-HO- + RH -. R- + H2O (1.4) (1.5) (1.6) (1.7) Termination + R O 2 + R-RO2 + R02-} inert products (1.8) (1.9) (1.10) Scheme 1.2 : F r e e - r a d i c a l chain mechanism for oxidation

F i g . 1 . 3 Relation between oxidation r a t e and hydroperoxide accumulation t"> .

With t h e a c c u m u l a t i o n o f h y d r o p e r o x i d e s , t h e r a t e o f b r a n c h i n g ( 1 . 4 - 1 . 7 ) i n c r e a s e s . Free r a d i c a l s formed i n ( 1 . 6 ) and ( 1 . 7 ) a r e new s o u r c e s o f h y d r o p e r o x i d e f o r m a t i o n ( 1 . 2 and 1 . 3 ) . T e r m i n a t i o n r e a c t i o n s end t h e c h a i n r e a c t i o n s , b e c a u s e i n e r t p r o d u c t s a r e formed. At some p o i n t t h e r a t e o f d e s t r u c t i o n o f t h e h y d r o p e r o x i d e s e x c e e d s t h e r a t e o f

formation, because termination reactions start dominating. Then the amount of hydroperoxides decreases. This is schematically shown in fig.1.3. The autocatalytic nature of the process is demonstrated by an

upward curve of the oxygen uptake after a so-called induction period

1.1.3 Photodegradat ion : Chemical Reaction Processes in PP

Degradation is studied by tracing the chemical reaction products formed, usually in liquid hydrocarbons or solutions. Infrared (IR) spectroscopy will be used in this research to monitor the increase of the amount of carbonyl and hydroperoxide groups. This method was chosen because it can easily be applied to solids. To illustrate the significance of the chemical groups studied, the photodegradation reactions causing their formation in PP will be described here briefly.

The review of Carlsson and Wiles [*-°l is referred to for more detailed

information. This covers the most important reactions in PP and is still considered to be up to datet°' l"> *-°l .

First of all the importance of hydroperoxides both as photoinitiators and in radical propagation is stressed. Hydroperoxide and peroxide groups are always present in PP, due to thermal degradation during processing or storage.

Because the tertiary hydrogen bond is easily dissociated, tertiary .peroxide

radicals—will_be-formed—i-n—reaction—1—1-Alkoxy radicals are formed by 1.4 or by photocleavage of the product formed by tertiary radical termination (1.10).

The reactions taking place after the formation of the alkoxy radicals give rise to the formation of compounds containing hydroxyl (~0H) and carbonyl (>C=0).

The 3-scission processes (reaction 1.11 and 1.12 in scheme 1.3) are examples. Reaction 1.12 is the most important because it leads to scission of the main chain of the molecule.

■ fc-CH2-C-CH2- + CH3- (1.11) II 0 CH3 I -CH2-C-CH2- " I 0-CH3

/

Scheme 1.3 : |S-sclssion processes in PP

Carbonyl compounds undergo photolysis by the so called Norrish

reactionsle-S-23] Moreover they absorb UV-light more efficiently than

hydroperoxides. This causes a continuing controversy about whether hydroperoxides or carbonyl groups are the main photoinitiators. Some carbonylgroups will usually be present in polyolefins as well as hydroperoxides.

The Norrish type I process (1.13, see scheme 1.4) for the type A ketone formed by (1.11) yields free radicals and causes chain scission, but has a very low quantum efficiency. It was observed by Garton et al.L"J that most ketonic species will be in the polymer backbone in PP and PE. In that case the two adjacent macroradicals formed in (1.13) will probably recombine before the initiation of oxidative chains can occur by O2 interception, so this process cannot be important in photoinitiation.

The Norrish type II process has a higher quantum efficiency but does not yield free radicals. According to Ranby and Rabekt^] (who at that time still accepted carbonyl groups to be important photoinitiators in PP) it occurs mainly in the type of keton formed in reaction (1.12), in which case it does not lead to additional chain scission. Because no free radicals are formed this reaction cannot initiate a chain reaction. Carbonyl compounds are therefore rather unimportant as photoinitiators in PP and contribute to the degradation process only after prolonged exposure when a large quantity of them has been formed.

In PE the discussion still goes on (compare[l°J a n d ! " ] ) . For PE the Norrish II process is dominant over type I

[25]

(1.14), and leads to end-vinyl groups. These are observed in PP as well but in much smaller quantitiesT1°]. 0 0 II h v || (I) ~CH2-C-CH2 - -> ~CH2-C- + -CH2- (1.13)

1

Scheme 1.4: Norrish type processes I (for the keton formed in reaction (1.11)) and II (for the keton formed in reaction (1.12)).

In summarising, the conclusion is justified that radicals formed by hydroperoxide decomposition are the most important photoinitiating

species in photooxidation of pp[24,39] -j^ie chemical reactions following

initiation, lead to the formation of ketonic groups, which are inefficient photoinitiators in PP.

1.1.4 Chain Scission and Cross Linking

Photodegradation alters the polymer's molecular structure and thereby changes its mechanical properties. The effects on the molecular structure are:

1) chain scission (breaking of the polymer molecule) 2) cross linking (connection of two polymer molecules)

The two effects are schematically drawn in fig.1.4. Chain scission results in a lower molecular weight causing a loss of ductility. This is

observed by a reduction of elongation and stress at rupture. The effects of extensive cross linking are a higher yield stress and Young's modulus[70•821 .

In most polymers both chain scission and cross linking occur, but the chain scission reaction usually prevails after some time, and the main effect of photodegradation in solid polymers is embrittlementt3] In polypropylene degraded by UV-light, chain scission is dominant. Cross linking is hardly observed.

Embrittlement occurs most rapid and severe in crystalline polymers, which is thought to be due to the occurrence of degradation in the

UV UV

Y

A

<

chain scission cross linking

Fig.1.4 Chain scission and cross linking

amorphous zones only (see also section 1.2.3). Localisation of degradation favours defect formation and consequent brittle fracture.

There are two factors which cause localisation in the amorphous zones in solid polymers. First of all, above the glass transition temperature (Tg j which is - -15°C for PP) the amorphous phase is

penetrable for oxygen while the crystalline phase is not, and secondly the photoinitiating impurities are rejected from the crystallites during crystallisation. So both the photoinitiating impurities and oxygen will be found in the amorphous phase.

From section 1.1.3 it will be clear that chain scission results mainly from the ^-scission process (reaction 1.12). Since carbonyl groups are formed in this process it is a reasonable assumption that there is a correlation between carbonyl formation, molecular mass reduction, and thereby mechanical deterioration. Therefore, carbonyl formation is often used as a chemical measure of the amount of degradation that has occurred.

1.1.5 Molecular Mass Reduction

The reduction of molecular mass due to degradation was observed for

many polymers I *■ • ■* i . In fact the reduction of the molecular mass is a

phenomenon that is more directly related to mechanical deterioration than carbonyl formation. Carbonyl formation can also occur in reactions that do not cause main chain scission ( e.g. reaction 1.11) and thus do not cause embrittlement. The molecular mass however, is more difficult to determine and therefore carbonyl-formation is more often reported. Studies which report both quantities can give an impression of the correlation between both phenomena.

Vink!.!" >'" J examined the relation between carbonyl formation, oxygen uptake and molecular mass reduction for a PP film (60 (jm thick).

A reduction of the viscosity average molecular weight Mv from 250,000 to

43,000 at the moment of embrittlement was reported.

Statistical theories were developed to relate the eventual molecular mass to the number of chain scissions

[4.31].

One of the more simple derivations is relation 1.15 :

5 - [(Mn)0/fïn>t]-l d - 1 5 )

where a = the average number of chain scissions per molecule at time t,

and Mn is the number average molecular weight. This formula is valid

only in conditions which do not apply to the crystalline polyolefins : random degradation, a small number of scissions and no volatilisation of products. Nevertheless it is often used as an indication of the amount of chain scission in crystalline polymersl*°.53-55]

To predict the molecular mass, a should be known and it would be

simple to obtain a^ if it would correlate to the amount of chemical reaction products formed. A thorough chemical analysis of the reaction products can yield information on the percentage of carbonyl forming processes leading to chain scission (e.g. 1.12), and the amount of carbonyl groups destroyed by other processes (e.g. 1.13).

Adams["-55] performed such an analysis for thermally degraded, process degraded and UV degraded PP sheets (20 mil. ,i.e.-500 (im thick). Products of chain scission reactions were followed by IR, and molecular weight by Gel Permeation Chromatography (GPC). At embrittlement of the

UV degraded sample, Mn decreased from 70,000 to 35,000 and the number of

chain scissions per molecule a-1 (Mv had decreased from 145,000 to

-65,000). In thermal oxidation the Mn decreased to -28,000 at

embrittlement, about the same level as after UV exposure. The number of chain scissions was a~2. The larger a at embrittlement can be ascribed

to the higher initial molecular weight (Mn=85,000).

But the carbonyl absorption was much larger in the case of thermal degradation than in UV degradation. In the first case two relevant reaction products per chain scission were found, and in the second case only one. This difference was explained by an oxygen diffusion limitation during UV degradation, which leads to hydrogen abstraction by one of every two radicals formed in the scission process.

The differences reported make clear that it is hazardous to convey results like this to other situations because many circumstances might influence the ratio of carbonyl content to chain scission. These may be the type of degradation process itself15*], exposure conditions,

stabilisation [*9T processing I55.57] orientation t5^l, morphology

[58]

stabilised films, and films with various stabilisers. In the non-stabilised film (100 |im thick), a reduction of Mn from 100,000 to about

10,000 was found (a-9), at a carbonyl content which was only half the level that Adamsl-^J observed in a similar case. However, there Mn had

only reduced from 70,000 to 10,000, and a-6.

It should be concluded that, though the molecular mass reduction is the link between chemical changes and mechanical deterioration, there is no simple way to quantitatively predict molecular mass reduction from chemical changes such as the formation of carbonyl groups, and molecular mass would have to be measured to obtain information about it. Results do seem to indicate that commercial PP with an initial Mv of -200,000 is

embrittled when M"v is reduced by UV degradation to about 40,000 to

60,000.

1.1.6 Effect of other Environmental Conditions

Apart from the UV-light intensity and spectral distribution, many other environmental conditions affect the degradation process. The reader is referred to refs.t-^"?] for their description. The most important ones are temperature and humidity. This research was limited to the degradation of specimens in the same artificial weathering equipment, so these were not varied.

1.1.7 Stabilisation

It is common practice to add degradation-preventing substances to polymers. These substances are called stabilisers. The two most

important classes are anti-oxidants (which stabilise against thermal

degradation during processing mainly) and UV-stabilisers (which are

specifically aimed against UV degradation).

The research on the understanding of the mode of operation of stabilisers is in full progress. Some successful new types have been developed recently. FormerlyIe-S- 3,47]j stabilisers were rigidly

classified by their mode of operation as:

- UV absorbers or screeners which prevent the absorbance of

the UV radiation by the polymer. They act by: 1) absorbing the harmful radiation and converting it to harmless forms of energy ; 2) reflecting the radiation and thereby preventing it from entering the polymer (as is done mainly by white pigments) ; 3) screening the deeper layers by a strong absorption in the surface (darker pigments, strong UV absorbers, surface coatings)

- excited state deactivators also called "quenchers", which

deactivate excited states

- peroxide decomposers , which convert hydroperoxides to

harmless non-free-radical products

- radical scavengers , which break the oxidative chain

either by removing an electron from an alkyl radical, or donating an electron to an alkylperoxyl radical

UV screening and quenching are specific for UV-stabilisers. Peroxide decomposers and radical scavengers are found among the anti-oxidants as

well. Anti-oxidants were not considered important in UV stabilisation at first because the known classes were found to be practically ineffective against UV-degradation. It is now recognised however, that a rigid classification is inapplicable and most successful photostabilisers act by enhanced simultaneous mechanisms.

Fig.1.5 illustrates the free radical degradation process of scheme 1.2 and the way it may be interrupted by stabilisers[10,11]

ROO-CB-D

Fig.1.5 Stabilising mechanisms : CB-A, chain breaking electron acceptor; CB-D, chain breaking electron donor; PD-S, peroxide decomposer (stoichiometric); PD-C peroxide decomposer (catalytic) ; UVA, UV absorber or screener ; MD, metal deactivator ; Q, excited state quencher 11">H J .

The former classification can be recognised in this figure (note that CB-A and CB-D are radical scavenging actions). The separate mechanisms are treated in the reviews of Gugumust^'land Allen et a l . l H ] . Scott [10] gives a number of commercial UV stabilisers and their mechanisms.

Some examples of frequently used UV stabilisers in polyolefins are [10,11,48,49] . derivatives of hydroxybenzophenon (UVA,CB-D), hydroxyphenylbenzotriazole (UVA, CB-D), hindered phenolic anti-oxidants

(CB-D) and nickel chelates ( UVA, CB-D, PD-C). A recent development which should be mentioned here are the very effective hindered amine light stabilisers (HALS). They prolong the half-time to as much as 40 times the original value[^°>^*'->2] .

In the current research, for reasons which will be explained later, a hydroxyphenylbenzotriazole derivative supplied by Cyba-Geigy (Tinuvin 326) was used to stabilise part of the specimens. In the case of a 100 Hm thick PP film, containing 1.0 wt/wt % of this stabiliser, the half-time was doubled compared to non-stabilised ppt^'j.

Conclusions

The photodegradation process is initiated by the absorption of light. Subsequent chemical reactions change the polymer structure. The main effect, influencing mechanical properties, is scission of the polymer backbone, which reduces molecular mass and causes embrittlement. There is no simple method to calculate the amount of molecular mass reduction based on reaction product formation only. Carbonyl formation, measured by IR-spectroscopy, yields a qualitative indication of the

amount of degradation that has occurred in polyolefins. When the Mv of

commercial PP degraded by UV light has decreased to about 40,000 to 60,000 the material will usually be severely embrittled.

1.2 Loss of Mechanical Properties

Introduction

In the following section literature on the decrease in mechanical properties of UV degraded polymers is summarised. The general trend is embrittlement of the material, which for example can be observed by the reduction of the stress and strain (or and er) at rupture in a tensile

test. The reduction of Gr and er of UV degraded polymers is rather

sudden. The time after which embrittlement occurs is often called the

mechanical induction time or embrittlement time analogous to the

chemical induction time.

The type of mechanical test and the specimen geometry have a marked effect on both the embrittlement time and the residual properties obtained.

Because there is a difference between thin films and thick specimens they are treated separately. In very thin films the degradation is not limited by oxygen diffusion

[5]

and the embrittlement time is short. Above a certain thickness., degradation takes longerle-g-3,32] because oxygen diffusion becomes a limitation. But for

very thick specimens, the embrittlement time does not increase further with thickness because degradation is limited to the surface only, and failure occurs through crack propagation.

The correlation between chemical and mechanical data can best be judged by fully degraded thin films: This subject is treated in section 1.2.1.

In thick specimens the degradation gradient through the specimen thickness results in a two layer system with an embrittled surface and a ductile core. Often crack formation occurs in the brittle surface. The formation and appearance of the cracks is treated in section 1.2.2.

The tensile behaviour of degraded thick specimens is described in section 1.2.3. Impact tests are treated in section 1.2.4 and flexural tests in section 1.2.5.

1.2.1 Thin films: Correlation between Chemical and Mechanical Test Results.

Embrittlement is caused by a lowering of the molecular mass through chain scission. The effect of molecular mass on mechanical properties is well studied. An empirical relation derived for the stress at rupture of amorphous polymers is (1.16)L^OJ_

Qr= A - ( —) (1.16)

nn

This relation gives a more or less steep, monotonous reduction of Qr

with Mn. However in UV degradation of crystalline polymers, this

relation will probably be invalid, because degradation is not homogeneous.

As was confirmed by V i n k t " ]j ^n thin films a bulk reaction is

expected i.e. there is no gradient of degradation through the film thickness. This makes them suitable for the study of the chemical

i n d u c t i o n t i m e i n s o l i d s , a n d the e f f e c t o n the m e c h a n i c a l p r o p e r t i e s . P P - i n d u c t i o n times r e p o r t e d v a r y c o n s i d e r a b l y d e p e n d i n g o n test c o n d i t i o n s (UV i n t e n s i t y , t e m p e r a t u r e , h u m i d i t y ) a n d s p e c i m e n type (processing, a d d i t i v e s t h i c k n e s s ) .

Pabiot a n d V e r d u t3 3] d e s c r i b e d the e f f e c t o n t e n s i l e p r o p e r t i e s b y

a so c a l l e d "failure e n v e l o p e " , f ( Or, £r, t ) . T h i s w a s shown t o b e the

original m e a n t e n s i l e c u r v e in the case o f P P . P E a n d P V C f i l m s , e v e n w h e n s t r e s s a n d e l o n g a t i o n a t r u p t u r e w e r e b e l o w the y i e l d p o i n t . T h e r e s u l t s , w h i c h a r e r e p r e s e n t a t i v e f o r d e g r a d i n g p o l y m e r films in g e n e r a l , a r e s c h e m a t i s e d in f i g . 1 . 6 .

800

£rl%) . tlh)

F i g . 1 . 6 L e f t : tensile curve a n d f a i l u r e e n v e l o p e , with f a i l u r e p o i n t s (in t h e left h a n d figure c l o s e d p o i n t s are n o n - d e g r a d e d s p e c i m e n s ) . R i g h t : a v e r a g e s of Br, £r a n d 0y

v e r s u s e x p o s u r e time. E x t r u d e d PP film 50 [im thick, 6 0 0 W A t l a s w e a t h e r o m e t e r l ^ J ] .

The

( er, Gr) .

left hand figure gives the coordinates of the rupture points The closed points here are undegraded specimens, whereas the open points are specimens, degraded for various times. The right hand

figures give the average strain er and stress or at rupture, and the

yield stress Gv vs. degradation time.

Most failure points are close to the original mean tensile curve (solid line). Considering the failure behaviour as a function of degradation time, it is observed that the rising portion of the cold

drawing part of the tensile curve is descended slowly. When Qr and er

fall into the horizontal part of the tensile curve, a considerable amount of scatter is observed. The right hand figures show that this part is crossed faster. The yield point finally is passed slowly, and eventually the specimen fails by brittle fracture. The steep part in the

reduction of er (right figure) coincides with the horizontal part of the

tensile curve where ds/de is very low. The authors suggest that "a localisation of chain breaks leading to defects on a supramolecular level" must be involved. If dG/d£ is small a strength reduction due to increasing defect concentration will sharply reduce the strain.

This also explains why usually the sudden reduction of Gr and er

does not coincide with a sudden reduction in Hn, and the generally poor

correlation between chemical and mechanical induction time. Some authors

report embrittlement well before the chemical induction period I3

This may mean that the very small amount of degradation occurring during the chemical induction period, is sufficient to reduce strength considerably. But more likely, chemical changes may have been averaged over an area in which local concentrations lead to defect formation. Possible causes for local concentrations are l^"J:

1 concentration of degradation in the amorphous domains, 2 stresses (residual or due to shrinkage) leading to crack

propagation and causing selective oxidation at the crack tips,

3 weak, oxygen sensitive regions such as spherulite boundaries,

4 a branching chain reaction causes a concentration of chain ruptures by itself.

Some remarks will be made on these possibilities :

ad 1 It was already mentioned in section 1.1.4, that degradation in crystalline polymers such as PP is restricted to the amorphous zones. The crystallites are impenetrable for oxygem'''and therefore remain unaffected. Moreover impurities and chain immobility cause local concentrations of degradation within the amorphous zones. In these zones "interlamellar molecules" are located, which (above T„) contribute largely to the toughness of semicrystalline polymers[^o] . Their scission is often suggested to cause rapid embrittlement.

Billingham and Calvertt^^J report that embrittlement occurred when the number average molecular mass was halved. They point out, that with

application of the "tie molecule concept", a reduction of Mn with a

factor five would be necessary (initial Hn-10^) to break sufficient tie

molecules to cause embrittlement. The "tie molecule" concept assumes that in the interlamellar areas, connecting the crystalline lamella, there is only a small amount of molecules which are connected the neighbouring lamella.

The tie-molecule concept is rather improbable in commercial polymers, which are rapidly cooled and where many interlamellar molecules are expectedI^1-^2] ^ good reason for selective oxidation of interlamellar molecules is absent in unstressed polymers. Therefore, explaining rapid embrittlement by the breaking of interlamellar molecules is "barely tenable" according to the authors. It should be noted that the concentration of degradation within the amorphous zones (caused by impurities and/or chain immobility) was not taken into account.

ad 2 Stress accelerates the degradation process. This effect would

occur in fully amorphous as well as crystalline polymers. Czernyl^JJ

showed that crack formation in thick parts is much accelerated by stress and facilitates oxygen diffusion at the crack tips. The question is whether tensile stresses occur in unstressed polymers. Residual stresses will be small in films, and in unstressed thick injection moulded specimens the surface will be in compression. However, in products with uneven wall-thicknesses and corners internal stresses may arise from uneven cooling.

In non-drawn films, stresses at spherulite boundaries occur from shrinkage during crystallisation and coolingl^'J,and in thick parts surface tensile stresses possibly build up through

"chemocrystallisation" (see also section 1.2.2). If a crack forms, stresses are high at the crack tip, and stress induced degradation will take place there, favouring further crack propagation.

ad 3 It was proven that in a crystalline polymer, impurities, irregularly polymerised material, additives and oxidised material 1*0] are rejected during crystallisation and driven to spherulite boundaries. This results in a concentration of chromophores at spherulite boundaries. Spherulite boundaries have been shown to be mechanically weak already (Friedricht->1]) and possibly they are also degradation sensitive.

ad 4 In solids the polymer radicals which carry the propagation reaction are relatively immobile compared to those in liquids and gases, in which the reaction mechanisms are usually studied. The reaction propagates through groups near the initiation site. This localises degradation and therefore reaction products will be found around photoinitiating impurities[22.28,37,40] ^n the amorphous zones.

The difference between the photodegradation of liquids and solids is a relatively new subject which will not be discussed further in this thesis. The reader is referred to recent surveys I " >38]

Another possible cause for the discrepancies between chemical and mechanical—induction—times—is—that—IR—spectroscopy^—usually-applied—to monitor chemical changes, is not sensitive enough.

Tüdös et al.[*"J measured carbonyl formation as well as molecular

mass (by GPC and by viscosimetry) . The 100 \im thick films that were

tested suffered a large reduction of molecular mass as well as toughness before carbonyl groups were detected by IR-spectroscopy.

When discussing films, most sources assume there is no degradation gradient over the film thickness. However, Carlsson and Wiles

[30] reported a higher level of UV-degradation in the outer 0.5 p of a 22

[im thick PP film. Averaging over the entire sample, mechanical

deterioration started while hardly any chemical changes were observed, but at the surface the chemical induction period already had passed. This was ascribed to chromophoric degradation products from thermal processing at the surface. So even in thin specimens a degradation gradient may occur, and techniques which can discriminate gradients over such small distances are usually not applied. It should be noted that Vinkt29] used the same techniques but the previous results were not

reproduced.

1.2.2 Thick Specimens : Surface Defects

Weathered thick specimens have a more or less homogenously degraded surface and an unaffected, ductile core. This is caused by factors which limit the depth of degradation^-. Fracture is often assumed to occur from surface defects^-S-".64-65]

1 This subject will be treated further in Chapter 5, concerning degradation depth.

Though many authors mention the formation of surface defects, a minute description of their formation and appearance is rarely given. Rapoport [■>"] reports cracks in thermally oxidised, thick (300 u m ) ,

oriented PP films, where degradation is limited by oxygen diffusion. The mechanism of crack formation is described as follows :"The surface is oxidised,- and shrinks. Because it is secured to the core, it cannot shrink freely, and mechanical stresses build up. In the amorphous zones some molecules are overstressed, and break easily under the action of oxygen as chemical agent. The elastic mechanical energy stored in the molecule is released, causing local heating, which accelerates the scission of other molecules. Where stresses are most concentrated, a surface microcrack will occur, causing a redistribution of stresses, followed by crack propagation." The cracks form perpendicular to the direction of molecule orientation , which was reported by others as well[3,66]) and they propagate to a depth of ca. 50 um. They are not straight but slightly curved (see fig.1.7.a).

Around the cracks the material was more severely degraded than in the rest of the surface. This may be due to local stresses as well as a better accessibility for oxygen. In fact the surface of the thick film is more quickly degraded than a thin film, taken from the surface, and degraded under the same conditions. This film did not develop cracks at the same degree of oxidation, which shows that these effects are at least partly due to the degradation gradient, causing local stresses.

Dolezelt^l shows cracks of ca. 300- 400 um deep in LDPE (weathered

two years in central Europe) and HDPE (weathered one year). The cracks are slightly curved as well, and the material around the cracks visually has a "different structure" and possibly is more degraded (see fig.l.7.b). In PP the cracks formed after only one month. Mechanical stresses are involved in crack formation, which is illustrated by the quicker formation of cracks when both specimen «nds are fixed to the back-ground.

a b

F i g . 1 . 7 . a surface cracks in PP films(300 pm t h i c k , t h e r m a l l y oxidised) [59], f i g . l . 7 . b surface cracks i n LDPE weathered for two years i n c e n t r a l Europe 1^1

These descriptions of crack formation suggest shrinkage of the

degraded polymer. Shrinkage, weight loss and an increase in density i s

often observed simultaneously in polyolefins Some authors ascribe

shrinkage and density increase to an increase of c r y s t a l l i n i t y due to

"chemocrystallisation,,[e-S-16>46,59-62] Billingham and Calvertt40]

describe this as the scission of chains which unravels entanglements and thus permits further crystallisation. Proof was obtained by IR spectrat16>°2land Small Angle X-ray scattering (SAXS)t60!.

However, other researchers could not detect an increase in crystallinity and ascribe density increase to the absorption of oxygen

by the amorphous phase. This causes weight increase, often observed at

the beginning of the oxidation process. Another cause may be the evaporation of low molecular products from the amorphous phase, which increases the percentage of crystalline material, and this would explain the shrinkage as well.

Whatever the cause, tensile stresses were shown to build up during

degradation by White and co-workers[44•63 ]| especially when the

reduction of the molecular weight was high. This occurred in amorphous (PS, PVC) as well as crystalline polymers (POM, PA 6 6 ) , and the surface usually deteriorated severely. Roughness increased, and sometimes fine surface cracks formed. After elongation up to rupture in a tensile test most specimens, including those which had no cracks before, showed cracks perpendicular to the elongation direction[44,45] Sometimes, in an extensively degraded specimen, the surface layer was reported to peel

off (Pet67!, PVC [45])

1.2.3 Thick Specimens : Tensile Tests

The most important features of the tensile behaviour of various weathered polymerste-6-3,5,44-45,66-70,72-73] a r e a s e v e r e reduction of

Er, a more modest reduction of or and thickness independence of the

embrittlement time. The effects in thick specimens are very much the same as in thin films, except that the changes usually take more time,

and er does not decrease to zero, but to a low stable end level.

The tensile curve is often descended as described in 1.2.1

(fig.1.6). In cases where 6r and sr are not very much affected the

material did not have a large cold drawing part in the tensile curve[44,69].

The yield strength and Young's modulus sometimes rise slightly in the beginning e.g. for PEl'0,82^ which is attributed to cross linking. At later stages Q „ and E may decrease, but the question is whether this is a bulk effect, or due to the reduction of the ligament by crack

formation in the surface. Finally Cy coincides with or, because the

sample fractures brittle before the onset of yield.

1.2.4 Thick Specimens : Impact Tests

Impact properties are more sensitive to degradation than tensile properties. High strain rate effectively raises the glass transition temperature and thus embrittles the polymer'. This makes the material behaviour more sensitive to surface defects (and possibly therefore to degradation effects) than a tensile test. Impact tests prove to be very

This is a generalisation. More careful discussion will follow in Chapter 6 on fracture toughness testing.

well suited to discriminate between materials for degradation resistance.

Impact tests are usually of the tensile, Izod or Charpy type. Impact bending tests are the most sensitive to degradation if the degraded side is in tension. Sometimes an instrumented drop weight impact test is applied, which is the most sensitive with the degraded side downwards!''1.

If the original material is very tough it will not fracture in an impact bending test. Either the supplied energy is too low, or the material tolerates large deformations without breaking. If the specimen is irreversibly deformed this can be regarded as a case of general yield. Usually after some degradation time a transition to fracture can be observed. In most cases however it is not the transition of gross yielding to fracture that is observed in impact, but a reduction of the fracture energy, as the fracture becomes more brittle.

Impact tests have been performed on degraded ABSt^°>7^.74-78]_

P C t7 2' " ] , P V C I7 3] , LDPEl8 0] A S A I7 6! , S A N I6 8] and P M M A I7 5! . The general

conclusion is that the fracture energy Wr reduces in a similar fashion

as Er in a tensile test, viz. an induction period, followed by a severe

reduction and a stable end level.

2.2.5 Thick Specimens : Flexural Properties

Flexural properties again are very sensitive to degradation, if the degraded side is tested in tension. The flexural strength Cv, reduces

more slowly than er in a tensile test, but quicker than Qj-I^T. Ruhnke

et al.["->'7°J developed a three point bending energy retention test,

where the fracture energy Wr(at 10 cm/min cross head rate) is measured.

They found a reduction of Wr similar to that of Er in a tensile test,

and Wr in impact, so again an initiation time, sudden reduction and

stable end level. Possibly because the effect of the high strain rate is absent, the fracture energy reduction is more modest than that of the impact energy.

Conclusions

The main effect of degradation on polymers is embrittlement. The tensile behaviour reflects this by a lowering of the stress and elongation at rupture. In this respect the behaviour of thick specimens is very similar to that of thin films. This is surprising because only a thin surface layer of the thick specimen is degraded, whereas the thin film is more homogenously degraded.

All information points in the direction that in thick specimens, surface cracks are the main cause of the fracture energy reduction. These cracks form during degradation or during the mechanical test, if the surface is sufficiently degraded.

Other tests, e.g. impact test and flexural energy tests show a reduction of the fracture energy as well. After a more or less steep reduction the fracture energy arrives at a stable end-level.