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The Influence of Crystalline Texture on the Tensile Properties of Natural Rubber. IAuthor(s): P. E. ReedReviewed work(s):Source: Proceedings of the Royal Society of London. Series A, Mathematical and PhysicalSciences, Vol. 338, No. 1615 (Jul. 16, 1974), pp. 459-478Published by: The Royal SocietyStable URL: http://www.jstor.org/stable/78593 .

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Proc. R. Soc. Lond. A. 338, 459-478 (1974)Printed in Great Britain

The influence of crystallinetexture on the tensilepropertiesof natural rubber. IBY P. E.REED

Department of Materials, Queen Mary College, London E. 1(Communicated by Sir Harry Melville, F.R.S. - Received 2 April 1973)

The crystalline morphologies that are attainable in samples of naturalrubber (n.r.), by extending the samples prior to crystallization, arereviewed. Specimens covering the full range of crystalline morphologiespossible have been prepared and tensile tested between - 120 and - 26?C.The tensile behaviour of crystalline samples is compared and contrastedwith that of oriented, but non-crystalline, identical natural rubber in thesame temperature range.It is found that the tensile behaviour of semi-crystalline n.r. is dominatedby the amorphous phase throughout the temperature range -120 to-26 OC.At temperatures above the glass transition temperature (Tg) ofthe amorphous phase, the crystalline phase acts mainly as a diluent of theamorphous phase. At temperatures below Tg,where the crystalline phaseis set in a glassy matrix, it is found that the crystalline morphology doessignificantly affect the tensile behaviour. Attempts are made to dif-ferentiate the effects of crystallinity, crystalline morphology and orienta-tion of the amorphous phase on the tensile properties of natural rubber.

1. INTRODUCTIONMany materials in common use are composite materials. In such composites asglass-reinforced plastics the constituents are deliberately selected and combinedto produce a material with particular properties. Studies may be performed eitherseparately on the individual constituents of these composites or on the compositeitself to verify theoretical predictions of the mechanical behaviour (Kelly 1966).There exists a further type of composite material in which the constituentsare chemically identical but exist in equilibrium with each other in differentforms. Semi-crystalline polymers containing a low proportion of crystalline phaseembedded in an amorphous matrix of the same polymer are such composites.No complete study of the variation in the mechanical properties of semi-crystalline polymers with crystallinity and crystalline morphology appears to havebeen attempted. In the case of polyethylene terephthalate, Dixon & Jackson (1968)have shown that large changes in crystallinity produce only marginal variations inyield strength, modulus and fracture stress at temperatures below the glass transi-tion temperature (Tg),while above Tgthese properties are highly dependent on thecrystallinity. Similar effects in raw rubber have been reported (Leitner I955). Theeffect of crystallinity on mechanical behaviour is thus dependent on the condition

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of the amorphous phase in which it is embedded. The importance of considering theamorphous phase in understanding the deformation of semi-crystalline polymershas been emphasized by Heffelfinger & Schmidt (I965).

Attempts have been made to investigate the effect of crystalline morphology onmechanical properties (Starkweather & Brooks I959; Pritchard, Dunn & KellyI964; Hammer, Koch & Witney 1959). Polymers in which the crystalline phase isalways spherulitic initially have been used in these studies, and the change inmechanical properties with spherulite size has been investigated. Unfortunatelychanges in spherulite size are accompanied by variations in crystallinity. Conse-quently it is difficult to separate the effects of these two parameters on the mechanicalproperties.A major problem in studying the mechanical behaviour of semi-crystallinepolymers is separating the effects of crystalline morphology, percentage crystal-linity, spherulite size and the contribution of the amorphous phase. Natural rubber(cis-polyisoprene) is a polymer which may be obtained in a semi-crystalline form.It offers certain advantages over other polymers for a study of the effect of crystallinetexture on the mechanical properties of semi-crystalline polymers. Natural rubbermay be rendered either amorphous or semi-crystalline, without modifying thestructure of the molecules, simply by storing the material at different temperatures.This feature permits the properties of the amorphous phase to be studied separatelyfrom those of the semi-crystalline composite. Natural rubber crystallizes relativelyslowly and the amount of the crystalline phase is easily controlled up to a maximumof about 30 % (Goppel I948; Nyburg 1954). The greatest advantage of using naturalrubber is the range of crystalline morphologies that may be generated by stretchingthe material to varying amounts before crystallization takes place (Andrews 1964).A further possibility is that the amorphous phase may be rendered either glassy orrubbery by testing the semi-crystalline material at temperatures either below orabove the glass transition temperature of the amorphous phase. Studies may thenbe performed with the crystallites set either in a glassy or an elastomeric matrix andin each case the crystallites will be bonded to the amorphous phase by primarybonds.

2. THE CRYSTALLINE MORPHOLOGY OF NATURAL RUBBERThe crystalline morphology produced in thin cast films of uncrosslinked naturalrubber has been previously studied (Andrews 1962, 1964). Thin films were stretchedby varying amounts at room temperature prior to crystallization and then held at- 26 ?Cfor crystallization to take place. The study of uncrosslinked natural rubberthin films showed that a wide range of crystalline morphologies were produceddepending on the 'pre-extension', i.e. the amount by which the films were extended

prior to crystallization. Unstrained films crystallized in a spherulitic manner, thespherulites consisting of crystalline islands interspersed with radiating crystallinefilaments and amorphous material. Small pre-extensions up to 50 % caused the

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Tensile propertiesof natural rubber. I

(a)

-? ?? 1'- :`';-'??3--*e----------C

n? r ? . rr ,?__----I I_T---rr--

(c)

(b)

. _ .

_ _

w -

-

..-- -__d-

- . - . .

(d)

, 1(e)

FIGURE1. Crystalline morphologies in thin films of natural rubber with increasing degreesof pre-extension: (a) 0 %, (b) 50%, (c) 100%, (d) 250%, (e) 500 % pre-extension. Thepre-extension direction is vertical.

spherulitic structure to become preferentially oriented with the filamentous com-ponent aligned perpendicularly to the direction of initial strain. When the pre-extension reached 100 %, the crystalline islands had entirely disappeared and onlycrystalline filaments running perpendicularly to the initial strain direction wereobserved. These filaments were termed a-filaments by Andrews. The a-filamentswere found to be arranged in ribbons or rows with their axes in the strain direction.Further increases in pre-extension beyond 100 % caused an increase in the numberand length of the rows of a-filaments together with a decrease in the length of thea-filaments themselves. In specimens subjected to very large pre-extensions, littleevidence of a-filaments was seen, but crystalline filaments running in the directionof strain were observed which were termed y-filaments. The density of the y-fila-ments in films pre-extended by 500 % was such that there was little room fora-filament growth. These five forms of crystalline morphology are sketched infigure 1.

Vol. 338. A.

461

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462 P. E. ReedMore recently the work on the crystallization in thin films of natural rubber hasbeen continued to study the effects of crosslinking on the process (Owen 1970).

Basically the same crystalline morphologies were found to exist as in the uncross-linked films at corresponding pre-extensions, except that crystallization rates,a-filament lengths and y-filament lengths were all reduced by crosslinking.The crystalline morphology in bulk samples of natural rubber is most likely athree-dimensional form of the morphologies observed in the thin films whencrystallized in a similar manner. Experimental studies to observe directly thecrystalline morphology in bulk samples, by taking thin sections from the crystallizedsamples for electron microscope investigation, have so far proved unsuccessful.A two-stage replica technique for electron microscope studies has been used to studyfracture surfaces of crystallized bulk samples of natural rubber, and revealsstructures in outline which are similar in size and shape to those observed in thethin films (Reed I970). The replicas show the outlines of spherulites in samplesreceiving no pre-extension prior to crystallization and the outline of ribbon shapessimilar in length, width and orientation to the rows of a-filaments observed inequivalently pre-extended thin films.Further evidence of the similarity of the crystalline morphology in thin films andbulk material is given by studies on polychloroprene (Andrews & Reeve 197I).The changes in crystalline morphology with pre-extension in polychloroprenefollow a similar pattern to those in natural rubber. X-ray studies on bulk samplesof polychloroprene, crystallized after varying degrees of pre-orientation, showedidentical orientation of the crystalline phase to that in thin films at equivalentpre-extensions. Hence crystalline lamellae in the bulk material correspond to the

TABLE 1pre-extension morphology expected morphologyratio, AS in thin films in bulk samples

1 spherulitic spherulitic1 < As < 2 preferentially oriented preferentially oriented

spherulitic spheruliticAs _ 2 a-filaments perpendicular to sandwich of lamellae plateletsinitial pre-extension direction and amorphous material withplane of platelets perpen-dicular to initial pre-extensiondirection2 < As < 4 row structure of a-filaments. lamellae platelet sandwicha-filament length decreasing arranged in stacks; lamellaeand row length increasing width decreasing and stackwith increasing A,; rows length increasing withoriented in pre-extension increasing A8; stacks orienteddirection in pre-extension direction4 < As < 7 as for 2 < As < 4 but with as for 2 < A, < 4 but withsteadily increasing y-filament steadily increasing y-filamentcontent content

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Tensile propertiesof natural rubber.Ia-filaments of the thin films at equivalent degrees of pre-extension. It has also beendemonstrated, using light scattering techniques, that samples of natural rubberstretched more than 300 % contain a light-scattering structure oriented in thedirection of extension (Yau & Stein 1968). This structure was identified with they-filaments observed in equivalently oriented thin films.The assumption that the crystalline morphology in the bulk samples is similarto that observed in thin films prepared under the same conditions of strain is madein the following interpretation of the mechanical test data. Such data as have beenobtained and discussed above support this assumption. The crystalline mor-phologies obtained with varying degrees of pre-extension (As) are summarized inthe above table.

3. EXPERIMENTAL(a) Specimen preparation

Dumb-bell shaped specimens were stamped from sheets of natural rubber 2 mmthick having the following composition and which had been cured at 130 ? for30min:parts bymass

smoked sheet 100stearic acid 1zinc oxide 5phenyl fl-naphthylamine 12-mercaptobenzthiazole 0.5sulphur 3

The specimens were mounted in suitable clamps for tensile testing and thenstretched at room temperature by the required amount to produce the desiredcrystalline morphology. The specimens were held at -26 ?C in their stretchedcondition. This temperature has been shown to be the optimum crystallizationtemperature for natural rubber (Gent I954). The time for maximum crystallinityto be achieved depends on the pre-extension. Unstretched specimens required20 days at -26 ?C for a maximum crystallinity of 24 % to be obtained, whilespecimens pre-extended by 200 % achieved their maximum crystallinity of 30 %in 2 days. Larger pre-extensions resulted in more rapid crystallization. Allspecimens were held at - 26 ?Cfor sufficient time to obtain maximum crystallinity.The tensile tests were performed in a special environmental chamber attached toan Instron Universal testing machine (Reed 1966). The chamber could be main-tained at any desired temperature in the range - 190 to 20 ?C. Since the crystallizedspecimens begin to melt at - 21 ?C,only 5 ?Cabove the crystallization temperature,a special insulated specimen transfer device was used to transport the specimensfrom the refrigerator to the pre-cooled testing chamber (Reed 1966).

29-2

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P. E. Reed(b) Tensile testing

A study was made of the effects of crystalline morphology and temperature oftesting on the tensile behaviour of semi-crystalline natural rubber. Specimens werepre-extended by amounts varying between 0(As = 1) and 600 % (As = 7) andcrystallized to give the complete range of morphologies shown in table 1. In allthat follows the pre-extension ratio can thus be regarded as an index of the crystallinemorphology present before testing commenced. Tensile tests were performed overthe temperature range - 120 to - 26 ?C for each morphology.Since the different crystalline morphologies were produced by varying the amountof orientation prior to crystallization, it is possible that the amorphous regionsremain oriented after crystallization. It is important to check whether observedmechanical behaviour is attributable to the presence of a particular crystallinemorphology or to the orientation of the amorphous regions. Hence all tensile testswere repeated using orientated natural rubber samples instead of semi-crystallinesamples. For these tests the rubber samples were stretched to the required amountat room temperature and immediately quenched to the test temperature, thusnot allowing time for crystallization to occur. No data for amorphous orientedspecimens are available for comparison with the semi-crystalline data when thepre-extension ratio exceeds 4 because natural rubber strain crystallizes rapidly atroom temperature when extended by more than 300 % (Nyburg I954; Mitchell &Meier I968).

All the tensile tests were performed at an extension rate of 2 cm min--.

4. RESULTS AND DISCUSSION(a) Types of mechanical tensile behaviour

Four different types of deformation were found within the experimental rangesof morphology and temperature studied, and may be typified by their respectiveload-extension curves shown in figure 2. The field of the present investigation maybe divided into four regions, each region being associated with a particular type ofdeformation. The first region is associated with type D curves and occurs withboth oriented amorphous and crystalline materials at temperatures above Tgwhenthe amorphous phase is rubbery. The three other regions exist below Tg. A brittleregion, associated with type A curves, occurs below Tgfor both types of materialat small pre-extensions. Ductile deformation, associated with type B curves,occurs for both types of material at intermediate pre-extensions below Tg. Withoriented amorphous samples the type B curve is associated with cold drawing, butwith crystalline samples stress whitening without cold drawing is observed. Thefourth region is also a ductile region, and may not be distinctly different from theother ductile region below Tg. It occurs at high pre-extension ratios in associationwith type C curves. The sharp load drop at yield of the type B curve is replacedby a change in slope on the type C curve.

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Tensile properties of natural rubber. I

strain-FIGURE 2. Typical stress-strain curves for the four different types of tensile behaviour ofnatural rubber in the temperature range - 120 to -26 ?C (arbitrary scales).

The iour types of behaviour are common to both oriented amorphous and semi-crystalline materials. This suggests that the basic deformation mode of the semi-crystalline material is that associated with the amorphous phase at all temperatures.However, comparison of the limits of temperature and pre-extension under whichthe four types of behaviour are obtained for the crystalline and oriented amorphousmaterials reveals significant differences. The inclusion of the crystalline phase innatural rubber apparently modifies, but does not dramatically vary, the mechanicalbehaviour of the amorphous material, except under test conditions close to transi-tions in deformation mode.

(b) Fracture stress and strainFigure 3 shows the variation of the true breaking stress (i.e. stress related to theactual cross-sectional area at rupture) with temperature and crystalline morphology.Reference to table 1 enables the pre-extension ratio (As)to be related to the crystal-line morphology present at the start of the tensile test. Figure 4 shows the totalextension to rupture (i.e. pre-extension plus further extension during tensile testing)as a function of the same parameters. Above - 70 ?Cthe fracture stress varies only

slightly over the complete range of morphology. The load-extension curves in thisregion were always of the type D form. Deviations from the nearly constant fracturestress occur only when As = 1 or As > 6. Thus, unless the crystalline morphology iseither spherulitic or has a high concentration of y-filaments, the fracture stressis not greatly affected by the crystalline morphology when the amorphous phase isrubbery. The extension to rupture (figure 4) follows the same pattern as the fracturestress above -70 ?C and exhibits a constant value for much of this range. Thisfurther endorses the conclusion that the crystalline morphology does not signifi-cantly affect the ultimate properties when the amorphous phase is rubbery.

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CD

cDI1

pre-extension ratioFIGURE 3. Variation of fracture true-stress withtemperature and crystalline morphology, indi-cated by the pre-extension ratio, for crystal-lized natural rubber.

FIGURE . Variation of the extreferred to the originalcrystallized specimen lengthand crystalline morphology,pre-extension ratio, for crrubber.

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Tensile properties of natural rubber. IA transition in the fracture stress occurs at about - 70 ?, which approximatesto the glass transition temperature of the amorphous phase. When As < 2.5 thistransition takes the form of a rapid decrease corresponding to a change to brittle

behaviour and type A load extension curves. The introduction of crystallinity inspherulitic form into a glassy amorphous matrix does not inhibit brittle failure inthe case of natural rubber.A transition to larger fracture stresses and elongation to rupture occurs at

temperatures below - 70 ?Cas the pre-extension ratio exceeds 2.5, and is associatedwith a change in the load-extension curve from type A to type B. This transition is

00.5 0

00/o,

10.3 2 .v

m ^o/ - I

o0.3

1 3 5 7 .9pre-extension ratio

FIGuURE . Change in fracture stress with pre-extension ratio at - 120 ?C for (o) semi-crystal-line natural rubber, ( x ) oriented natural rubber and (v) acrylonitrile-butadiene rubber.

coincident with the change from an aligned spherulitic morphology to the fullydeveloped lamellar sandwich structure with the lamellae set normal to the strainaxis initially. The transition may alternatively be related to the increasing orienta-tion of the amorphous phase, since a similar transition from brittle to ductilebehaviour is observed with the oriented amorphous samples, although in the lattercase the transition occurs at lower pre-extension ratios.The fracture stresses of semi-crystalline and oriented non-crystalline specimensfor tensile tests performed at - 120o0 are compared in figure 5. The data for non-crystalline natural rubber are not given beyond As = 3.5, since at this pre-extensionstrain crystallization occurs at room temperature. However, data for acrylonitrile

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468 P. E. Reedbutadiene rubber (abr), which is a non-crystallizing synthetic rubber, are includedin figure 5. The graph shows that the fracture stress of crystalline material increasesfivefold in the range 1 < As < 5 while the fracture stresses of the oriented materialsexhibit only a two fold increase. Hence orientation of the amorphous phase, althoughit increases the fracture stress, cannot alone account for the large rise in fracturestress of the semi-crystalline material. Figure 5 also shows that the fracture stressof the semi-crystalline spherulitic material (As= 1) at -120?C, which contains24 % crystallinity, is the same as that of the glassy amorphous material alone.Hence the dramatic rise in fracture stress of the semi-crystalline material withchanging morphology cannot be explained in terms of changes in the crystallinecontent, which are in any case small over the range tested. The fivefold increase infracture stress of the crystalline material is thus attributed mainly to changes in thecrystalline morphology. At

- 120?C and at high pre-extensions (large y-filamentcontent) the fracture stress decreases abruptly and brittle behaviour again obtains.The fracture stress values when As > 6 appear to be an extrapolation of the brittlestrength values when As < 2.5.Further evidence that the change in mechanical behaviour at As = 2.5 in semi-crystalline samples is associated with morphological changes is provided by thefracture surfaces of specimens tested below - 70 ?C. When As < 2.5 the fracturesurface is a single plane transverse to the draw direction and typical of brittlefracture. When As > 2.5 a fibrous fracture surface is obtained and bunches offibrous material are found in the vicinity of the fractured specimen. Such fibrousmaterial is consistent with the fracture path going round the stacks of lamellae ory-filaments.

(c) The region of rubber-like behaviourStress-strain curves for the complete range of crystalline morphologies studiedfrom tensile tests performed at -40 ?Care shown in figure 6. The inset on figure 6shows the initial form of the curves to a larger scale. Corresponding curves for theoriented non-crystalline specimens are shown in figure 7. These latter curves are also

reproduced dotted in figure 6 for ease of comparison. The starting extension ratiofor each curve indicates the pre-extension ratio or initial morphology.Figure 7 shows that the stress-strain curves of the oriented specimens fall on onesingle curve regardless of pre-extension ratio providing As < 5. Such behaviour isentirely consistent with rubber elasticity theory which states that the stress is afunction of state, and independent of the path taken to arrive at that state. Speci-mens cooled to - 40 ?0 and then extended, or extended at room temperature andthen cooled, exhibit similar stresses at a given extension. The deviation from thecommon curve in figure 7 when As > 5 can be attributed to stress relaxation asso-ciated with crystallization (Gent 1954) during preparation for testing. No commoncurve is apparent for the semi-crystalline samples (figure 6), but rather the curvesare of similar form for all morphologies and displaced along the extension ratio axis.Comparison of the curves for the crystalline and oriented samples shows thatthe major effects of introducing the crystalline phase are to increase the rate of

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extension ratioFIGURE 6. Simple elongation stress-strain curvesfor semi-crystallized natural rubber at -40 ?C.Extension ratio values plotted include both ex-tension during tensile testing and pre-extension

prior to crystallization. Dotted curves indicateequivalent data for oriented amorphous speci-mens. Insert shows start of some curves to alarger scale.

extensFIGURE 7. Simple elongfor oriented amorph- 40 ?C.Extension raboth test extension a

temperature. Pre-exte*, A = 1.8; A, A8 = 5.5; o, A8= 7.

*+*Pt-FCDmq2caQtl

h3

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strain hardening and reduce the overall extension to fracture. Curves shown infigure 6 are for the maximum crystalline content of approximately 30 %. Lesscrystallinity results in curves intermediate between the solid curves in figure 6 forcrystalline and the dashed curves for oriented amorphous samples. On initialexamination of the results it appears that the crystalline phase may merely serveas a filler in the amorphous matrix. The modulus of the crystalline phase mustexceed that of the amorphous phase when the latter is rubber-like, and, from con-siderations of simple composite theory (Holliday I966), the modulus of the com-posite will increase with crystalline content. The enhancement of the modulusshould depend only on the volume fraction of crystalline phase present accordingto simple composite theory. Since the crystalline content is approximately constantfor all samples tested, the stress-strain curves for all semi-crystalline samples maybe expected to fall on a common curve. Clearly no such common curve exists for thecrystalline samples in figure 6, even after allowance is made for the pre-extension.It is thus considered that both the crystalline content and to a lesser extent thecrystalline morphology affect the form of the stress-strain curve at temperaturesabove the Tg of the amorphous phase, although the ultimate strength does notapparently vary with the morphology.Variations in the modulus in small strain tensile tests have been attributed tochanges in the aspect ratio (i.e. ratio of length to width) of the crystalline phase(Halpin & Kardos I972), the modulus increasing with aspect ratio. When 2 < As < 5the aspect ratio of an individual lamella within a stack of lamellae is considered todecrease with increasing pre-extension as outlined in table 1. Such a reduction inaspect ratio should result in a decrease in modulus according to Halpin & Kardosrather than the increase observed. The stiffening effect may result from the stacksof lamellae acting in concert to serve as fibrillar reinforcement, since the aspect ratioof a stack of lamellae increases with pre-orientation in contrast to that for anindividual lamella. Alternatively the increasing modulus, or strain hardening rate,with pre-orientation may result from orientation of the amorphous phase. Evidenceof residual orientation is shown by the stress required to maintain the specimens atthe pre-orientation extension after crystallization. Stress-strain curves in figure 6do not commence at zero stress when As > 2.5. The initial stress on each curveindicates the stress required to maintain the pre-extension after crystallization, andincreases with pre-extension. Although these stresses are small, they correspond tosignificant extension in the amorphous phase if the same residual stresses are appliedto the amorphous phase alone. The increase in residual stress with orientationsuggests increasing residual orientation in the amorphous phase after crystallization.Such orientation can contribute to the increasing strain hardening rate with pre-extension as will be discussed in a subsequent paper.

(d) Behaviour belowTgThree regions of distinctly different behaviour are observed below Tg for bothoriented amorphous and semi-crystalline materials. Figures 8 and 9 show, for the

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Tensile properties of natural rubber. I

-0.24Qe

exte.nA4^y^si*L^ ^ ^"1202tio ' ^*:8 0 -1001ot20temperatureFIGURE8. Variationof brittle fracturestress and yield stress with temperatureand initialorientation for oriented amorphousnatural rubber. Brittle fracture and yield areindicatedby shadedand plain surfacesrespectively.

_0.4

-0.3

-02

.0.1

temperaturel CFIGURE. Variationof brittle fracturestressandyield stresswith temperatureandcrystallinemorphology, indicated by pre-extension ratio, for semi-crystalline natural rubber.Brittle fracture and yield are indicated by shaded and plain surfaces respectively.amorphous and semi-crystalline materials respectively, the variation in brittlefracture stress or yield stress with temperature and degree of pre-extension. Theregion of brittle fracture is shown by the shaded region and the fracture stress isplotted. The plain surface indicates the variation of the yield stress where ductilebehaviour obtains. A trough appears in each of the yield surfaces in figures 8 and 9,with the lowest point at approximately As = 3 in figure 8 and As = 4 in figure 9.At pre-extensions below these values ductile deformation is associated with type B

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curve (figure 2), and a load drop at yield is observed. At greater pre-extension valuesthan the low point of the trough, ductile deformation is associated with type Ccurves. In the latter case the yield stress is taken as the stress at the abrupt changein slope of the stress-strain curve.Each type of behaviour occurs with both oriented amorphous and semi-crystallinematerials, but under slightly different conditions of temperature and pre-orientation.Hence the fundamental behaviour must again be attributed to the deformation ofthe amorphous phase.(i) Region of brittlebehaviour

The shaded areas in figures 8 and 9 show the regions where brittle fracture obtainsfor oriented amorphous and semi-crystalline materials respectively. Comparison ofthe two shaded areas shows that (a) the fracture stresses of the semi-crystallinesamples exceed those of the oriented amorphous samples at all values except atAs = 1, (b) the region of brittle behaviour extends to higher As values in the caseof the semi-crystalline samples at temperatures below - 90 ?C,(c) brittle behaviouroccurs immediately below Tgwhen As = 1 (unoriented) in the case of the amorphousmaterial, but does not occur until about 20 ?C below Tg in the case of the semi-crystalline material containing spherulites.The identical fracture stresses of the amorphous and semi-crystalline materialswhen As = 1 show that the inclusion of a crystalline phase does not automaticallyincrease the brittle fracture strength. The enhanced brittle fracture stress in thesemi-crystalline samples when As > 1 must be attributable to the arrangement ofthe crystalline phase (i.e. crystalline morphology), rather than the mere presence ofcrystallinity. This view is supported by further tests using a series of specimenswhich contained mainly parallel lamellae (figure 1c), alined at various angles to thetensile test strain axis initially. The greatest enhancement of the brittle fracturestress occurred when the lamellae were normal to the test direction. Samples inwhich the lamellae were alined in the direction of tensile testing showed no enhance-ment of the brittle fracture stress over that of the unoriented amorphous material.It is not clear to what extent the lamellae width affects the strengthening mechanismin the brittle behaviour region, since a transition to ductile deformation occurs withincreased pre-orientation before any significant change in the lamellae width isattained.(ii) Brittle-ductile transition

Both oriented amorphous and semi-crystalline materials yield and undergosubsequent plastic deformation at temperatures below Tg, provided the pre-extension is sufficiently large. The minimum pre-extension required to cause ductilebehaviour increases with decrease in temperature as shown in figures 8 and 9.The similarity in behaviour of the oriented amorphous and semi-crystallinesamples suggests that yield is associated with the amorphous phase, since this iscommon to both cases. Since the brittle-ductile transition occurs over a narrow

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Tensile propertiesof natural rubber.Irange of pre-orientation, any effect of crystalline morphology on the transition isnot apparent. The presence of the crystalline phase inhibits yielding, increasing thepre-extension required to change the behaviour from brittle to ductile by approxi-mately 100 % at each test temperature below - 90 ?C. Results from the orientedamorphous material alone show that, at any particular temperature, yield occurswhen the orientation of the amorphous phase achieves a critical value. If yieldinitiates in the amorphous phase at a critical value, the shift of the brittle-ductiletransition to higher pre-extension in the case of the semi-crystalline samples suggeststhat the orientation in the amorphous phase of the semi-crystalline material is lessthan in the oriented amorphous counterpart at any given pre-extension.At temperatures between - 90 and - 70?0 the semi-crystalline material is ductileover the whole range of pre-orientation, although the oriented amorphous materialis brittle for small degrees of pre-orientation. Non-homogeneous deformation occurswhen a spherulitic structure is present. Instead of one load decrease at yield on theload-extension curve prior to the propagation of a single yield band, several loaddecreases are observed as extension proceeds, producing a saw-tooth form of load-extension curve. Each load decrease is associated with the formation of a new yieldband in the specimen. Any attempt at interpretation of the yield mechanismoccurring in the semi-crystalline samples containing a spherulitic structure is mademore difficult by the non-homogeneous deformation.

0.05

0.03 I.crystalline oriented- 0.03

' 0.01-0 i J I I, I1 3 5 7 9

extensionratioFIGURE 10. Stress-strain for semi-crystalline and oriented amorphous natural rubber at- 80 ?C,for specimens pre-extended by 80 % at room temperature prior to crystallizationand/or tensile testing. Extensions plotted include the pre-extension.

If yield initiates in the amorphous phase, subsequent deformation is certainlymodified by the presence of a crystalline phase. The stress-strain curves for orientedamorphous and semi-crystalline samples, having pre-extension ratios of 1.8, arecompared in figure 10. Extension at constant load after initial yield is associatedwith cold drawing in the case of oriented amorphous samples. Termination of thedeformation at constant load at an extension ratio of 4, regardless of pre-orientation,suggests that the amorphous phase seeks to achieve a natural draw ratio of 4 onyielding. The semi-crystalline samples exhibit strain hardening immediately after

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yield, contrasting with the oriented amorphous behaviour. The specimens deformuniformly throughout their length, at least macroscopically, at the test strain ratesadopted. Stress whitening occurs as deformation proceeds, possibly due to theformation of voids as a result of incompatibility of the amorphous and crystallinephases during straining. Mechanisms of yielding in semi-crystalline polymers havebeen discussed by other researchers (Peterlin 197I; Keller & Pope I97 ).

(e) Elastic modulusThe initial tangent modulus is used to define the elastic modulus at very smallstrains, since non-Hookean behaviour occurs at temperatures above Tg.The varia-tion of the moduli of the semi-crystalline and oriented amorphous materials with

temperature and pre-orientation are shown in figures 11 and 12 respectively.Comparison of these figures shows that the changes in moduli with temperature andpre-orientation are similar for both types of material. The usual large increase inthe modulus occurs with both types of material at the glass transition of theamorphous phase. The similarity of figures 11 and 12 shows that the amorphousphase dominates the small strain elastic behaviour of the semi-crystalline material.There are, however, small variations between the two figures which are significant.At temperatures above Tgthe modulus of both types of material increases withpre-extension. The modulus of the oriented amorphous material is always slightlygreater than that of the semi-crystalline material when As > 2.4. This contradictspreviously reported data (Leitner 1955; Starkweather & Brooks I959; Halpin &Kardos I972) showing that the introduction of a crystalline phase increases themodulus. In all these previous cases, however, crystallinity was introduced withoutinitially orienting the amorphous phase, in contrast to the presently reported work.

Halpin & Kardos have used composite theory to explain the effect of increasingcrystallinity on the modulus. When the modulus of the introduced phase exceedsthat of the matrix in which it is set, the modulus of the composite exceeds that ofthe matrix, provided the modulus of the matrix is not changed by the addition ofthe new phase. In the present case the modulus of the semi-crystalline sample iscompared with that of the oriented amorphous material at equivalent pre-exten-sions. If the properties of the amorphous phase are different in the two samples, themoduli of the respective amorphous phases will differ. If composite theory is invokedin the present case, the results may be explained by suggesting that the modulus ofthe amorphous phase in the semi-crystalline sample is less than that of the orientedamorphous sample alone when As > 2.4, and the inclusion of the crystalline phasedoes not compensate for this difference. Such differences in moduli of the amorphousphases could be due to variations in extension, since the modulus of the amorphousmaterial increases with strain. This interpretation is consistent with the earliersuggestion that the orientation of the amorphous phase of the semi-crystallinematerial is less than in the oriented amorphous counterpart at equal pre-extensions.The discrepancy between the present work and that reported by other researchers,on the effect of increasing crystallinity on the modulus, is thus attributed to

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Tensile propertiesof natural rubber.I

f ?r

E.OS0

Pre-extensOn rato 240 eloCFIGURE11. Initial tangent modulus plotted against temperature and crystalline morphology,indicated by pre-extension ratio, for semi-crystalline natural rubber.

pre-extension ratio t,eFIGURE 12. Initial tangent modulus plotted against temperature andpre-extension for oriented amorphous natural rubber.

variations in orientation of the amorphous phase resulting from the method ofintroducing the crystallinity.At temperatures below Tg,a small increase in orientation from the non-orientedstate results in a twofold increase in modulus for both types of material. This againemphasizes the dominating role of the amorphous phase. At pre-extension ratiosgreater than 2 the modulus of the oriented amorphous material was found to beessentially constant for all pre-extensions, although considerable scatter of the data

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P. E. Reedwas observed in this region due to the experimental method used. It is apparent,however, that the modulus of the semi-crystalline material exceeds that of theoriented amorphous equivalent at A,s 2 and A, > 6. At values intermediatebetween these values the modulus of both types of material is similar. The crystallinemorphology thus appears to affect the elastic modulus at A, - 2 and As > 6 attemperatures below Tg.At large values of pre-extension the modulus of the semi-crystalline materialagain rises rapidly. The crystalline morphology at these values of pre-extensionis predominantly that of y-filaments alined in the direction of testing. Fracturesurfaces from such specimens are highly fibrous. Hence the y-filaments serve asfibre reinforcing for the glassy amorphous phase, and higher elastic moduli result.

5. CONCLUSIONThe mechanical behaviour of a semi-crystalline polymer is determined by thebehaviour of the crystalline and amorphous phases acting in concert. Generally itis difficult to separate the contribution of each phase to the observed behaviour ofthe polymer. Interpretation of the tensile behaviour of natural rubber has been

helped in the present investigation by the ability to study the behaviour of theoriented amorphous material separately from that of the semi-crystalline com-posite. A systematic experimental study has been made of the effect of crystallinemorphology on the tensile behaviour of natural rubber over a wide temperaturerange. This behaviour has been compared with that of oriented amorphous naturalrubber throughout the same range. Four regions of distinctly different tensilebehaviour have been identified within the range of the investigation, which arecommon to both semi-crystalline and oriented amorphous materials. This similarityin behaviour shows that, for the material tested, the amorphous phase dominatesthe tensile behaviour at temperatures both above and below the Tgof the amorphousphase: the maximum crystallinity possible in natural rubber is approximately 30 %.Although the stress-strain curves for the oriented amorphous and semi-crystallinematerials are similar within any one particular region of tensile behaviour, smalldifferences are significant in understanding the role of the crystalline morphology.At temperatures above Tgthe inclusion of the crystalline phase increases the rateof strain hardening and decreases the extension to rupture. It appears that theseeffects are related mainly to the crystallinity present and are not greatly affected bythe form of crystalline morphology. The tensile strength above Tg, calculated astrue stress, is not modified by the inclusion of the crystalline phase. The sameproperty expressed as engineering stress shows an apparent increase in strengthwith crystallinity, since variations in strain at rupture are not then taken intoconsideration. The crystalline phase thus appears to act mainly as a diluent of theamorphous phase when the latter is an elastomer. Small variations in the form of thestress-strain curves for the semi-crystalline samples may be attributable todifferences in crystalline morphology or may be due to variations in the orientation

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Tensile propertiesof natural rubber.Iin the amorphous phase resulting from the method of forming the different crystal-line morphologies.At temperatures below Tg, the inclusion of the crystalline phase causes a shiftin the ductile-brittle transition to higher temperatures at any given value of pre-extension. Consideration of the data for the oriented amorphous material aloneshows that the transition is related to a minimum critical orientation, whichincreases with decreasing temperature. It is considered that the formation of acrystalline phase in the equivalently oriented sample causes a relaxation in theamorphous phase, and results in the shift in brittle-ductile transition to highertemperatures for the crystalline samples. Below Tg t is found that a fivefold increasein fracture stress occurs within the range of crystalline morphologies studied. Suchan increase cannot be explained on the basis of variations in crystalline content or interms of changes in amorphous phase orientation. This dramatic increase in tensilestrength is thus attributed to changes in the crystalline morphology. The maximumstrength is attained with a morphology which is thought to be stacks of lamellaearranged in a crystalline-amorphous sandwich and set in an amorphous matrix,the stacks being alined in the test direction. The lamellae stacks appear to act asfibre reinforcement since the fracture surfaces of these stronger samples is fibrillar.

Comparison of the tensile behaviour of oriented amorphous and semi-crystallinenatural rubber has shown that the amorphous phase dominates the behaviour in thetemperature range - 120 to - 26 ?C.The crystalline phase appears to act mainly asa diluent of the amorphous phase when the latter is an elastomer; variations incrystalline morphology having little effect on the tensile behaviour. When theamorphous matrix is glassy, however, it appears that both crystalline morphologyand crystallinity affect the tensile strength, elastic modulus and the form of thestress-strain curve.

The author is indebted to Professor E. H. Andrews for suggesting the topic ofinvestigation and many useful discussions. The author also thanks the NaturalRubber Producers Research Association for providing the samples for theinvestigation.

REFERENCESAndrews, E. H. I962 Proc. R. Soc. Lond. A 270, 232.Andrews, E. H. 1964 Proc. R. Soc. Lond. A 277, 562.Andrews, E. H. & Reeve, B. I97I J. Mat. Sci. 6, 547.Dixon, E. R. & Jackson, J. B. I968 J. Mat. Sci. 3, 464.Gent, A. N. I954 Trans. Farad Soc. 50, 521.Goppel, J. M. I948 Rubber Technology Conference, London.Hammer, C. F., Koch, T. A. & Whitney, J. F. 1959 J. appl. Poly. Sci. 1, 169.Halpin, J. C. & Kardos, J. L. 1972 J. appl. Phys. 43, 2235.Heffelfinger, C. J.. & Schmidt, P. G. I965 J. appl. Poly. Sci. 9, 2661.Holliday, L. 1966 Composite materials. Amsterdam: Elsevier.Keller, A. & Pope, D. 197I J. Mat. Sci. 6, 453.Kelly, A. I966 Strong solids. Oxford University Press.Leitner, M. 1955 Trans. Farad. Soc. 51, 1015.

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478 P. E. ReedMitchell, J. C. & Meier, D. J. I968 J. Poly. Sci. A-2, 6, 1689.Nyburg, S. C. 1954 Br. J. appl. Phys. 5, 321.Owen, P. I970 Crystallization in thin films of natural rubber. Ph.D. Thesis, University ofLondon.Peterlin, A. 197I J. Mat. Sci. 6, 490.Pritchard, R., Dunn, T. & Kelly, P. 1964 J. appl. Poly. Sci. 8, 1751.Reed, P. E. I966 J. Mat. Sci. 1, 91.Reed, P. E. I970 The influence of crystalline texture on the tensile properties of naturalrubber. Ph.D. Thesis, University of London.Starkweather, H. W. & Brooks, R. E. I959 J. appl. Poly. Sci. 1, 236.Yau, W. & Stein, R. S. I968 J. Poly. Sci. A-2, 6, 1.