POLYMER WEATHERING: MECHANISMS OF DEGRADATION AND FAILURE

POLYMER WEATHERING: MECHANISMS OF DEGRADATION AND FAILURE J R White* School of Chemical Engineering and Advanced Materials, Newcastle University, Newcastle upon Tyne, NE1 7RU, UK jim.white@ncl.ac.uk The origins of the deterioration of the engineering properties of polymers in outdoor service are discussed and links are made between events at the molecular level and the engineering failure of polymers caused by weathering. The aim of this contribution is to discuss ways in which characteristics such as chain scission rate can be measured and used to explain changes (such as changes in crystallinity) occurring at different scales and so to account for changes in the sensitivity to fracture of components of macroscopic dimensions. It is shown that a combination of characterization techniques such as FTIR, GPC, XRD, DSC, SEM and residual stress analysis can be used to build up a fairly comprehensive understanding of failure mechanisms in weathered polymers. Introduction The intention of this review is to show how study of several different aspects of degradation are required to develop an understanding of polymer weathering. It is important to make measurements of changes occurring at the molecular level and of macroscopic behaviour in order to be able to assess the weatherability of polymers. It is necessary to determine changes occurring at several intermediate length scales to understand how changes occurring at the molecular level produce the observed changes in engineering properties. Because the length of this paper is strictly limited it is impossible to provide a comprehensive review of this topic. Therefore I use examples drawn mainly from my own research: I trust that the many authors who have made fine contributions in this field will understand if I cannot make reference to their work. A further rationalisation is that the review concentrates on photo-oxidation: although it is generally accepted that exposure to solar UV radiation is the prime cause of degradation in outdoor exposure, other elements can have a significant influence. Aspects of outdoor exposure such as humidity, rainfall, the presence of gaseous pollutants can all affect polymer degradation but are touched upon only briefly or not at all. Thus the survey begins with an example of a technique (Dynamic CO 2 monitoring) used to probe the reaction chemistry of samples under UV exposure. The most important consequence of (photo-)chemical is change in the macromolecular architecture, usually in the form of chain scission or crosslinking. A powerful way to obtain insights into these changes is the measurement of the molecular weight distribution and this is discussed next. When chain scission occurs, molecular segments are often released from entanglements and if this occurs in a semi-crystalline polymer the segments that are so liberated may contribute to secondary crystallization and the crystallinity of the

material increases ( chemi-crystallization ). Quite apart from providing a characteristic for following the development of degradation, crystallinity measurement provides a direct link to engineering properties. This is because the crystal phase is very much stiffer than the amorphous phase and has quite different thermal properties. Thus chemi-crystallization causes a large change in the engineeirng properties and this important topic is introduced here. When UV irradiation is strong, the changes caused by photo-oxidation that are described above occur preferentially near to the surface. This is because the oxidation process is so rapid that oxygen is consumed near the surface before it can diffuse very far into the interior of the polymer. Reaction in the interior is therefore very slow, apart from the time immediately after the illumination is switched on, when there will be oxygen available that diffused into the polymer during the preceding dark period. This phenomenon is known as oxygen diffusion limited reaction and gives rise to a distinctive depth profile of degradation beneath the exposed surface. In the case of a semi-crystalline polymer the chemi-crystallization occurs primarily near to the surface. The crystal phase is denser than the amorphous phase with most polymers and chemi-crystallization with be accompanied by surface shrinkage that is resisted by the material in the interior that remains largely unchanged. Therefore the surface may go into a state of tensile residual stress, a very detrimental development. Residual stress analysis is a rather neglected aspect of polymer characterization but is dealt with here. Finally a section is included on mechanical testing. Even if a polymer is not used specifically as a load bearing component it must retain its mechanical integrity in order to perform whatever function it is designed for, and an assessment of the mechanical degradation is a necessary part of any investigation into its weatherability. The results are interpreted in terms of the observations made by the other techniques and it is attempted to make links between the observations made at different levels. Dynamic CO 2 monitoring Many of the techniques used to monitor polymer photodegradation require exposure times of the order of several days to several weeks in order to observe significant changes. In contrast, the dynamic CO 2 method can yield meaningful results in tests lasting of the order of 3 hours. It can be used to determine the relative sensitivity of different polymers to photodegradation or to compare the effect of different additives on the susceptibility of a polymer to photodegradation. The method was developed initially to monitor paint photodegradation 1. The sample is placed in a glass cell and exposed to UV provided by a 150 W xenon tube via a flexible light guide and a CaF 2 window 2,3. Filters are used to control the spectral distribution of the illumination. The cell is placed in a FTIR instrument so that the interrogating IR beam passes through the cell parallel to the sample surface,

perpendicular to the UV illumination. The IR beam passes into and out of the cell via CaF 2 windows and analyses the gaseous reaction products emitted from the sample. The reaction takes place in a chosen gas that is flushed through the cell for an hour prior to sealing and commencing the experiment. Details of the experimental procedure are given elsewhere 2,3. The progress of photodegradation is followed by monitoring the build up of carbon dioxide within the cell using the IR band centred at 2360 cm 1. The majority of experiments conducted with this method to date have used samples based on polyethylene (PE) or polypropylene (PP), with a variety of additives, including a range of different TiO 2 grades 2,4. With PE compounds, CO 2 generation was detected within 10 minutes of switching on the UV illumination, as soon as was practically possible, whereas with PP a distinct induction period was observed 4. Complementary information on the progress of oxidation was obtained using conventional UV exposure followed by FTIR analysis, focusing mainly on carbonyl group development, and the results were explained in terms of the detailed reaction mechanisms believed to apply to these polymers 4-7. Development of carbonyl absorption in PP is slower than for PE, and with PP, oxidation is dominated initially by formation of perhydroxyl, not carbonyl, groups. The slow development of PP carbonyl groups appeared to correspond to the ~30 minutes induction time for CO 2 photogeneration obtained with PP, but not with PE. Pre-exposure of either polymer to UV increased subsequent photogeneration of CO 2, suggesting that CO 2 is formed from carbonyl groups and that, for PP, CO 2 production is delayed because, unlike PE, the initial oxidation products are perhydroxyls. The elimination of a CO 2 induction time for PP films in which carbonyl groups had been induced by pre-exposure to UV supported this interpretation. Although the evidence for reaction mechanism provided by the CO 2 method is somewhat circumstantial, it is provides a way to compare polymer induction behaviour in a relatively short test period. Recording of spectra is normally continued for a further hour after the irradiation is switched off. With some polymers, CO 2 emission stopped almost immediately but for others, such as poly(vinyl chloride) 3, emission continued, presumably the result of continuing oxidation promoted by radicals that had accumulated during exposure or because CO 2 generated deep within the sample required time to diffuse to the surface and escape. Again, the CO 2 method provides information that relates to the complex mechanisms that are involved in photodegradation. Molecular weight distribution The molecular weight of a polymer is often determined using gel permeation chromatography (GPC) and this is the technique used by the author and co-workers. It is common to generate molecular weight averages (number average, M n, or weight average, M w ) but this method also

provides the molecular weight distribution (MWD) and this contains more information. Polyolefins such as PE and PP require high temperature GPC, which is not very common but has been used for most of the work reported here 8. When a polymer is photodegraded the molecular weight averages normally fall, the result of chain scission. Crosslinking is a competing process with some polymers, however, and the molecular weight averages rise if it dominates. In most reported studies chain scission dominates and the whole of the MWD shifts towards lower molecular weights, sometimes with little change in the shape of the distribution. If significant crosslinking occurs, a high molecular weight tail sometimes develops and indicates that some molecules are larger than those in the original population, prior to the photodegradation 9,10. When both scission and crosslinking occur, molecular weight averages no longer reflect the extent of degradation that has occurred. Shyichuk has developed a method of analysing MWDs to separate the different molecular degradation mechanisms 11. It involves comparing experimental MWDs with MWDs generated using Monte Carlo computer-aided modification of the starting MWD, assuming scission and crosslinking are both random events (Molecular Weight Distribution Computer Analysis: MWDCA). This procedure has been used to examine the nature of photodegradation in polystyrene and several polyolefins 10,12-17, and the effect of the inclusion of stabilizer and/or pigment 16,18. It has revealed very strong differences between the degradation mechanisms displayed by different polyolefins. For example, in a low density polyethylene (LDPE) scission concentrations rose at a fairly uniform rate during the early stages of UV exposure whereas for a PP homopolymer the scission rate rose very steeply near the exposed surface, which could be an indication of the presence of an induction phenomenon or of auto acceleration, which occurs when the products of reaction are themselves pro-degradants 14,17. In the interior, scission and crosslink rates were smaller and fell with extended exposure 17. This is attributed to oxygen diffusion limited reaction; the fall in reaction rates probably occurred because, in the early stages, some dissolved oxygen was present and/or because the increase in reaction rate at the surface reduces the amount of oxygen diffusing into the interior. When a photostabilizer was present the molecular weight averages fell fairly uniformly throughout the sample depth 19. This is because the much-reduced oxidation allowed oxygen to diffuse into the interior so that it was available to participate in reactions there. More detailed analysis using MWDCA revealed that the scission rate increased with exposure time, showing that, even at the much-reduced reaction rates obtained when stabilizer was present, autoacceleration could still occur. When samples were stressed during UV exposure it was found that molecular weight averages fell more rapidly when tensile stress was applied whereas compressive stress caused a modest reduction in the fall in M n or M 20-22 w. Application of MWDCA to the data confirmed that scission rates increased and also showed that

the scission/crosslink ratio increased when tensile stress was applied. Although it is tempting to suggest that tensile stress will assist separation of broken molecular fragments and discourage recombination, this reflected a general observation that at higher overall reaction rates the scission /crosslink ratio increased. Crystallinity measurement Semi-crystalline polymers often have a lamellar structure in which thin ribbon-like crystals are constructed from molecule segments. The molecules fold at the lamellar surface ( fold surface ) but do not necessarily re-enter at the adjacent site. Molecules pass through the crystal lamellae and the surrounding amorphous phase, providing strong adhesion between the two phases. Adjacent lamellae are connected by tie molecules that enter more than one of the crystal lamellae, and pass through the intervening amorphous material, and by molecular entanglements within the amorphous phase. The crystalline and amorphous phases have very different properties and the mechanical properties depend strongly on the fraction of material in each phase. Thus it is of importance to have methods to measure the (fractional) crystallinity. This can be done using X-ray diffraction (XRD) or differential scanning calorimetry (DSC). XRD crystallinity measurement X-ray crystallinity is normally measured using Bragg-diffraction peaks obtained with a diffractometer. The amorphous phase scatters X-rays into all directions and the crystalline and amorphous diffraction components are normally separated relatively easily. The X-ray intensities in the crystalline and amorphous parts of the scattering envelope are in proportion to the mass fraction of the respective phases 23. There are geometric factors to be applied (Lorentz factor etc) 23. DSC crystallinity measurement DSC crystallinity is obtained from the crystal melting endotherm obtained during a DSC heating run. In order to estimate the fraction of material that was in crystal form prior to melting it is required to know the melting enthalpy for a fully crystalline sample. Very few polymers form defect-free crystals with no amorphous phase, so this quantity cannot be measured directly and it must be deduced from measurements made on semi-crystalline samples. For samples cooled very rapidly when crystallized from the melt state (e.g. material in the skin of an injection moulding, frozen by contact with the cold mould wall), some crystallization may occur during the DSC heating run before reaching the crystal melting temperature as molecules become mobilized at the elevated temperature ( cold crystallization ). This is a potential source of error in the measurements that are meant to represent the crystallinity of the material prior to running the DSC test.

This method can be used to determine whether chemi-crystallization has been provoked by UV exposure 24. If the sample is cooled down in a controlled manner in the DSC cell, then reheated, a second heating thermogram can be obtained. The memory of the original processing history and of any chemi-crystallization is erased by the first melting in the DSC and the second run gives information characteristic of the material only. In the case of a photo-degraded polymer the crystallinity may increase, if the main result was chain scission, producing a sample in which molecular entanglements are less likely to inhibit crystallization, or it may decrease, because molecular defects such as carbonyl groups have formed and do not fit into the crystal lattice 25. Comparison of XRD and DSC crystallinity measurements Both XRD and DSC crystallinity measurements assume a simple two-phase crystal-amorphous morphology. However, the material immediately adjacent to the crystal surface will have a structure that is intermediate between a fully ordered crystal and the random chain structure associated with the amorphous phase. This is sometimes referred to as the inter-phase and it is expected that there will be some contribution from this material to the crystalline diffraction peaks, producing some broadening, and distinct from the amorphous scattering. The inter-phase will influence crystal melting and will have a secondary effect on the size of the melting endotherm. It is unlikely that the effect of the inter-phase on the crystallinity measurements will be the same for the two techniques and it is reasonable to accept values that are not exactly equal. A further difference between the two methods concerns molecule segments freed from entanglements by photo-initiated chain scission. They may require elevated temperature to become sufficiently mobilized to display secondary crystallization: they may therefore provide enhancement of the crystal signal obtained in DSC (if the cold crystallization contribution cannot be separated) but not in XRD. Furthermore, the X-ray diffractometer obtains data from near the surface, limited by the X-ray absorption characteristics of the material, whereas DSC samples are extracted by microtoming or milling a layer of finite thickness (usually 0.1 mm in the studies reported here) from a chosen depth within the sample. Although the X-ray measurement can be performed on a surface at a chosen depth from the exposed surface by microtoming or milling away material, the sample depth is not the same, making exact comparison difficult if there is a strong depth profile present. Reasonably good agreement between XRD and DSC crystallinities were obtained from a series of PP copolymer samples at different depths from the exposed surfaces after different UV exposures 26. Clear evidence was obtained for chemi-crystallization. Such contributions are expected to diminish at higher exposures because the molecule segments progressively acquire molecular defects. This will also affect cold crystallization and is therefore expected to cause DSC measurements to level off quicker than XRD measurements if the cold crystallization contribution is not subtracted, and this was observed 26.

Residual stresses Residual stresses form in thermoplastics as the result of the large thermal gradients that are present during moulding operations. For example, in injection moulding, hot polymer melt is injected into a mould and the material in contact with the cold mould wall solidifies rapidly, forming a skin with dimensions matching the cavity. The material in the interior then cools slowly; thermal shrinkage is prevented by the solidified skin and tensile shrinkage stresses develop 27,28. They are opposed by compressive reaction stresses that form near the surface. The compressive stresses inhibit cracking from surface flaws and are generally regarded as beneficial (as in toughened glass). Regrettably, outdoor exposure has been found to cause the stress distribution to reverse, giving tensile stresses near the surface 29,30. This is extremely detrimental because tensile stresses are then present in a region that has been embrittled by photodegradation and which can easily form flaws that act as stress concentrators from which cracks may be driven by the tensile stresses. In the case of semi-crystalline polymers, tensile stresses develop near the exposed surface because chemi-crystallization occurs preferentially there: crystallinity measurements made through the depth of PP bars in the as-moulded state and after 9 weeks laboratory UV exposure confirmed this. Shrinkage occurs because the crystalline material is denser than the amorphous material from which it forms. Calculations of the shrinkage caused by the change in crystallinity were converted into stress and this accounted for the majority of the change in residual stress observed 31,32. Similar stress reversal has been observed in non-crystalline polymers 29. This is believed to be the result of accelerated physical ageing in which the molecules in the amorphous phase rearrange to produce more compact, denser structures 33. The accelerated ageing occurs preferentially near to the surface for the same reason as chemicrystallization, through the release of molecular entanglements that inhibit the rearrangements, as the result of photo-oxidative chain scission. Laboratory tests have shown that the application of a temperature gradient can also reverse the sense of the as-moulded residual stress distribution; exposure to sunlight may cause such a temperature gradient to develop and is an alternative explanation for the results obtained with samples exposed outdoors 34. Mechanical testing Weathering causes deterioration of mechanical properties, sometimes very quickly, and mechanical testing is an essential part of assessing weatherability. Although loss of toughness is usually the reason why weathered polymers fail, impact tests are not especially common in studies of polymer weatherability. This is partly because the pendulum-based tests (Izod, Charpy) are not very suitable. This is because they use a notch for a crack starter that is introduced artificially. This immediately

poses the problem of when to introduce the notch. If it is introduced after outdoor exposure it is likely to cut through the most heavily degraded material at the surface and cracking will commence within material that has not degraded very much. Conversely if the notch is cut before exposure it will expose material with different morphology to that in which photo-oxidation normally occurs, given that injection mouldings and other fabricated polymers have a steep morphology gradient at the surface. Although this objection does not apply to drop-weight impact tests on plaques, it is found that conventional uniaxial mechanical tests dominate the mechanical assessment of weathered polymers. The most sensitive parameter is then the elongation to break. Ductile polymers such as PP that cold-draw (to several 100%) during a conventional tensile test become brittle and break after strains less than 10% after UV exposures of a few weeks unless they are stabilised. This is a good indication of toughness and normally a good guide to service behaviour. It is advisable to conduct post-test fractography to determine the fracture mechanism. This is especially important when laboratory artificial weathering trials are conducted to speed up the assessment of weatherability. Degradation is accelerated by a variety of means including: use of high UV intensity; 24 hour per day exposure; elevated temperature; temperature cycling; water spray and/or high humidity. If the fracture mechanisms in tensile tests on samples weathered outdoors and in the laboratory, respectively, are different, accelerated laboratory testing cannot be used to predict service performance or to compare weatherability of different polymers. An example where SEM fractography confirmed that artificial weathering provoked the same type of mechanical failure can be seen on comparing figure 7(a), reference 35 (natural weathering of PP) and figure 3, reference 36 (artificial weathering). In both cases photo-oxidation caused formation of a brittle surface layer, <0.5 mm deep, in which cracks formed and propagated easily; the crack then arrested when it encountered the less degraded material in the interior and a patch of ductile fracture occurred before the crack was long enough to cause rapid fracture across the rest of the section. Another phenomenon that has been observed in studies of the mechanical properties of photodegraded polymers is recovery 37,38 in which the strength of a polymer is sometimes found to be greater after a longer period of UV exposure than after a shorter one. This happens as the result of the following sequence: (i) UV exposure causes a brittle surface layer to form that is sufficiently deep to cause net section brittle failure under tensile loading and a large fall in strength; (ii) further UV exposure causes the surface layer to become weaker, allowing multiple cracks to form (mutually unloading one another) and preventing stress transfer into the relatively undegraded interior so that the surface cracks no longer propagate once they have penetrated through the surface layer, and the strength recovers somewhat. Sometimes the surface layer becomes so degraded that it flakes off altogether. More examples and discussion are given elsewhere 36,39-42. It should be noted

that the recovery is observed in carefully controlled laboratory experiments but is of no practical use, since failure is likely to have occurred in service before the recovery phenomenon has a chance to occur. The important practical lesson from these observations is that the interval between tests on exposed samples must be kept fairly small so that the poorest properties, prior to recovery, are not missed, since this would give an over optimistic impression of the weatherability of the material. Conclusions Complex chemical reactions occur during weathering of polymers and lead to a variety of molecular, morphological and physical changes that combine to control the engineering behaviour. In some polymers, changes in molecular size may lead to chemi-crystallization. The secondary crystallization is likely to be located near to the surface in a thick-sectioned sample and causes tensile residual stresses to form. The residual stresses are likely to cause reduction in the strength of the material. Surface embrittlement leads to net-section brittle failure but further degradation of the surface layer may allow recovery in carefully conducted laboratory tests; this should influence the way test programmes are designed but has no other practical value. Acknowledgements This paper is based largely on research conducted by the author and the many collaborators, postdocs and students whose names are found in the reference list that follows. Their friendship and dedication to their research has sustained me throughout my career. References 1. P.A. Christensen; A. Dilks; T.A. Egerton; J. Temperley J.Mater.Sci. 1999, 34, 5689. 2. C. Jin; P.A. Christensen; T.A. Egerton; E.J. Lawson; J.R. White, J.R. Polym.Degrad.Stab. 2006, 91, 1086. 3. C. Jin; P.A. Christensen; T.A. Egerton; J.R. White. Mater.Sci.Tech. 2006, 22, 908. 4. S.S. Fernando; P.A. Christensen; P.A. Egerton; J.R. White Polym.Degrad. Stab. in press 5. P. Delprat; X. Duteurtre; J.-L. Gardette Polym.Degrad.Stab. 1995, 50, 1. 6. P. Gijsman; J. Hennekens; D. Tummers Polym.Degrad.Stab.1993, 39, 225. 7. F. Gugumus Makromol.Chem.Makromol.Symp. 1989, 27, 25. 8. B. O'Donnell; J.R. White; S.R. Holding J.Appl.Polym.Sci.1994, 52, 1607. 9. B. O'Donnell; J.R.White J.Mater.Sci. 1994, 29, 3955. 10. A.V. Shyichuk; J.R. White J.Appl.Polym.Sci. 2000, 77, 3015. 11. A.V. Shyichuk; V.S. Lutsjak Eur Polym J. 1995, 31, 631.

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