Orientation of the a- and c-modification of elastic polypropylene at uniaxial stretching A. Boger a, * , B. Heise a,1 , C. Troll b,2 , O. Marti a,3 , B. Rieger b,2 a Abteilung Experimentelle Physik, University of Ulm, Albert-Einstein-Allee 11, D-89069 Ulm, Germany b Abteilung Anorganische Chemie II, University of Ulm, Albert-Einstein-Allee 11, D-89069 Ulm, Germany Received 7 December 2006; received in revised form 10 April 2007; accepted 18 May 2007 Available online 2 June 2007 Abstract New metallocene catalysts applied to propylene polymerization expand the range of properties of polypropylene (PP), resulting in semi-crystalline materials having crystallinities below 60% up to X-ray amorphous highly elastic ones. To date the origin of the unique elastic mechanical behavior of such low crystalline PP is not completely understood. Therefore, the microscopic orientation of those PPs due to uniaxial stretching was investigated using wide-(WAXS) and small-angle X-ray scattering (SAXS). The aim of this study was to correlate these orientations or changes in the developed fiber tex- tures with the macroscopic stress–strain behavior. This includes efforts to come closer to the main question of the nature of the physical cross-links in these not chemical cross-linked homopolymers, which is the reason for the high elastic behavior. Therefore, high molecular weight metallocene PPs showing different crystallinities (0–36%) were stretched to several elon- gations and the structural changes during the deformation were recorded by X-ray scattering. Stress–strain measurements show the great potential of these PPs as a thermoplastic rubber material. For quantitative analysis and discussion of the polymer chain orientations, the orientation functions were calculated. Correlations between the orientation functions and the stress–strain curves allow an interpretation of the macroscopic behavior on a microscopic scale. A higher cross-linking density in elongated samples indicates that the network, which is responsible for the elasticity, mainly built up by strain- induced morphology changes and chain orientations. Ó 2007 Elsevier Ltd. All rights reserved. Keywords: Elastic polypropylene; Low crystalline; Metallocene catalysts; Orientation; X-ray scattering; Stretching 1. Introduction In the last decade, metallocene catalysts gained rapidly increasing interest in academic research as well as in industrial development [1]. Previously the design of newly designed metallocene catalysts has led to a series of new polypropylene microstruc- tures, like highly syndiotactic materials, hemi-iso- tactic chains [2,3,18] and atactic–isotactic block 0014-3057/$ - see front matter Ó 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.eurpolymj.2007.05.031 * Corresponding author. Present address: R&D Biomaterials, Synthes GmbH, Eimattstrasse 3, CH-4436 Oberdorf, Switzer- land. Tel.: +41 61 965 64 09; fax: +41 61 965 66 04. E-mail address: [email protected](A. Boger). 1 Tel.: +49 731 50 23008. 2 Tel.: +49 731 50 23038. 3 Tel.: +49 731 50 23010. European Polymer Journal 43 (2007) 3573–3586 www.elsevier.com/locate/europolj EUROPEAN POLYMER JOURNAL
Transcript
EUROPEAN
European Polymer Journal 43 (2007) 3573–3586
www.elsevier.com/locate/europolj
POLYMERJOURNAL
Orientation of the a- and c-modification of elasticpolypropylene at uniaxial stretching
A. Boger a,*, B. Heise a,1, C. Troll b,2, O. Marti a,3, B. Rieger b,2
a Abteilung Experimentelle Physik, University of Ulm, Albert-Einstein-Allee 11, D-89069 Ulm, Germanyb Abteilung Anorganische Chemie II, University of Ulm, Albert-Einstein-Allee 11, D-89069 Ulm, Germany
Received 7 December 2006; received in revised form 10 April 2007; accepted 18 May 2007Available online 2 June 2007
Abstract
New metallocene catalysts applied to propylene polymerization expand the range of properties of polypropylene (PP),resulting in semi-crystalline materials having crystallinities below 60% up to X-ray amorphous highly elastic ones. To datethe origin of the unique elastic mechanical behavior of such low crystalline PP is not completely understood. Therefore, themicroscopic orientation of those PPs due to uniaxial stretching was investigated using wide-(WAXS) and small-angleX-ray scattering (SAXS). The aim of this study was to correlate these orientations or changes in the developed fiber tex-tures with the macroscopic stress–strain behavior. This includes efforts to come closer to the main question of the nature ofthe physical cross-links in these not chemical cross-linked homopolymers, which is the reason for the high elastic behavior.Therefore, high molecular weight metallocene PPs showing different crystallinities (0–36%) were stretched to several elon-gations and the structural changes during the deformation were recorded by X-ray scattering. Stress–strain measurementsshow the great potential of these PPs as a thermoplastic rubber material. For quantitative analysis and discussion of thepolymer chain orientations, the orientation functions were calculated. Correlations between the orientation functions andthe stress–strain curves allow an interpretation of the macroscopic behavior on a microscopic scale. A higher cross-linkingdensity in elongated samples indicates that the network, which is responsible for the elasticity, mainly built up by strain-induced morphology changes and chain orientations.� 2007 Elsevier Ltd. All rights reserved.
In the last decade, metallocene catalysts gainedrapidly increasing interest in academic research aswell as in industrial development [1]. Previouslythe design of newly designed metallocene catalystshas led to a series of new polypropylene microstruc-tures, like highly syndiotactic materials, hemi-iso-tactic chains [2,3,18] and atactic–isotactic block
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structures [4]. This expanded the range of propertiesfor iPP and sPP enormously [3,5–8,18]. That cata-lyst technology creates also an elastomer using pro-pylene (PP) as a feedstock [1]. Elastomers provideelasticity to products such as athletic clothing, con-sumer packaging and impact-resistant automotiveand industrial components. The new catalysts createthe first family of low crystalline, high molecularweight PPs with high elastomeric properties. Theelasticity is due to new polymer microstructures incontrast to conventional Ziegler–Natta polypropyl-enes [2,8–11]. Metallocene based polymerizationwas described by Brintzinger and coworkers[12,13] who used stereo rigid bridged C2-symmetricligand moieties around the active zirconium center.Depending on the catalysts symmetry and polymer-ization conditions, a great variety of PPs with differ-ent stereo-regularities and molecular weightdistributions, resulting in different crystallinities,can be polymerized with different physical structuresand properties [3,18].
Thermoplastic Elastomers (TPEs) combine themechanical properties of vulcanized elastomers(rubber) with the process ability of thermoplastics.The property of TPE to be repeatedly melt able isdue to the lack of chemical cross-links comparedto conventional rubber material. There are severaloptions to create physical networks resulting in elas-tic polymers. Either the material exhibits elasto-meric properties through copolymerization orblending of different thermoplastic materials. Inthe case of copolymerization, rigid segments formdomains, which act as physical cross-links. A com-mon way to get PP with elastic properties is to blendhighly isotactic PP with low isotactic PP or withPolyethylene [14,15]. The blending could also bedone with PP samples having different isotacticityresulting from one PP material after fractionation[16], as first described by Natta [17]. Another possi-bility is to produce amorphous block copolymers,where one block sequence has a higher glass transi-tion than the temperature range of application andlowers the other block sequence. The hard elementsact as cross-links, building up the necessary networkin the soft matrix, which allows the low stress defor-mation. Those hard elements can also be stablesmall crystals, which are isotropic distributed in ablock copolymer [18] or in a high molecular weightstereoblock PP described by Collette et al. [19–21].The structures of these TPEs result in a low distor-tion temperature and high compression set at roomtemperature compared to cross-linked elastomer.
This leads to a more economical process due toshorter cycle times with the advantage of recyclables[14].
Thus, a broad variety of elastomeric polypropyl-enes has been synthesized [3,4,6,15,18,22–30], butthe relation of their morphology and mechanicalbehavior is not fully investigated. Studies performedon stereoblock isotactic polypropylenes reveal, thatthe polymers crystallize mainly as short individ-ual lamellae embedded in an amorphous matrix[20,21,26,31–33]. Schonherr et al. [31–33] revealedfor such polypropylenes with mmmm contentsbelow 30% morphologies comparable to classicalsemi-crystalline polymers, like lamellae, crosshatch-ing, hedrites, and spherulites. Since the amount andthe length of the isotactic sequences are known toinfluence the crystallization behavior [34–36], thecorrelation between the chain microstructure orien-tation due to stretching and the resulting macro-scopic behavior observed are of particular interestto understand how and to what extent these chainorientations can explain the resulting mechanicalproperties. Results of the mechanical studies ofsemi-crystalline polymers reported in literatureclearly show, that properties like tensile strength,elongation at break or reversible deformation areassociated with rearrangement and deformation ofthe crystalline and amorphous domains [37]. Basedon many studies on highly oriented semi-crystallinepolymers the microscopic process of tensile defor-mation was proposed to proceed within severalregimes [38–48].
The present study was performed using highmolecular weight PPs with different stereo-regulari-ties resulting in different crystallinities. The elasticPPs were synthesized with new metallocene cata-lysts, which were activated with MAO (MethylAlu-minumOxane) or borate [3,18,36]. These catalystsshow a high activity towards the polymerization ofpropylene and lead to significant high molecularweight products [3,18,36]. The amount of stereoerrors in the chains of these PPs influences thedegree of crystallinity, resulting in semi-crystallinepolymers and even nearly amorphous ones withhigh elasticity [36,49,50]. The microscopic orienta-tion during uniaxial stretching of the disordereda/c net-planes [50] was determined using X-ray scat-tering techniques and correlated to the macroscopicmechanical behavior. There are only a few studiesconcerning the mechanical behavior and chain ori-entation during stretching of such low crystallineiPP [6,51] with special stereo-regularities and high
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molecular weights [3,18]. Investigated materialsherein are not studied in material science for a longtime as iso- [52] or syndiotactic [7] PPs having tactic-ities >55% and melting transitions above 80 �C[7,53–55]. Materials having elasticity with reversiblerecovery above 80% as described herein present anew class of PP. A really worthwhile detailed anal-ysis of the possible and preferred orientations ofcrystals that develop during stretching metallo-cene-made iPP samples was modelled and presentedin the work from Auriemma and De Rosa [51].
The present work is an extension of previousinvestigations on the same elastic metallocene poly-propylene samples [50]. Herein a deeper understand-ing of the unique mechanical behavior of thosematerials is presented.
2. Experimental section
Four metallocene polypropylene samples showingdifferent mechanical properties and crystallinities wereinvestigated in the present study. Methods and resultsfor polymerization, chemical material characteriza-tion, sample preparation and tensile tests presentedherein are basically the same like in the previous study[50], added here for completeness and allow discussionusing tensile curves of the samples.
2.1. Polymerization, chemical material
characterization and sample preparation
Polymerization experiments [56] were per-formed in toluene solution and in liquid propyleneat temperatures between 283 and 323 K in a 0.5 LV4A Buchi steel reactor equipped with internaltemperature control and an efficient cooling sys-tem (Tp = const. = 275 K). The activators MAOor trityl-tetrakis(pentafluorphenyl)borate wereadded to a solution of catalyst in the polymeriza-tion medium at the desired temperature. Water
Table 1Crystallinity, chemical characterization of the investigated polypropyle
Detailed analysis of the microstructures data is given in [36] with the corand P4 herein, respectively.
impurities were scavenged by triisobutyl aluminumbefore metallocene addition. The polymerizationreaction was quenched by injection of methanol;excess monomer was vented off, polymer was pre-cipitated out of toluene solution in methanol andall polymers were dried in vacuum at 363 K.
13C NMR spectra were recorded on a BrukerAMX 500 spectrometer (C2D2Cl4, 353 K, with aminimum of 10 000 scans, 2.2 s delay time) and ana-lyzed by known methods [57] to determine the pen-tade distribution. Molecular weights and molecularweight distributions were measured by gel perme-ation chromatography (Waters, alliance GPC2000, 418 K in 1,2,4-trichlorobenzene) universal topolystyrene and relative to polypropylene standards(Table 1). For more details, the micro structuraldata of the samples are presented in the work fromHild et al. [36]. Identification of the samples investi-gated herein to the mentioned manuscript [36] isgiven in the legend of Table 1.
Investigations presented were performed usingraw material as received from polymerization andmelt pressed films. For the latter purpose, 10 g ofthe polymer materials were cut into granules, placedin between two glass plates (distance of 1 mm) on aheating stage with load control and covered withPTFE foils. The polymer was heated under constantload (2 kN) to 420 K. When the temperature wasreached and granules were molten, the load wasincreased to 5 kN. After leaving the sample for30 min in this state, the load was increased to20 kN. After another 30 min, the pressed film wascooled to room temperature with a cooling rate of1.5 K/min. The circular films had dimensions of1 mm in thickness and 100 mm in diameter. For ten-sile testing and X-ray scattering investigations rect-angular strips with nominal dimensions of 70 mm inlength and 10 mm in width were stamped out of thefilm. Areas on the film showing inhomogeneitieswere excluded from the tests.
ne samples
2/% Density ± 0.01,g/cm3
Polymerization-medium/temperature
Centermetal/cocatalyst
0.9 toluene/303 K Zr/MAO0.87 toluene/323 K Zr/MAO0.85 toluene/308 K Zr/MAO0.85 liquid PP/293 K Hf/Borate
responding material names A5, A3 and B1 for the materials P2, P3
Fig. 1. Sketch for the explanations of the read outs from 2dscattering images for analysis.
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2.2. Tensile testing
To investigate the mechanical properties undertension, rectangular strips were submitted to tensiletests at room temperature, using a commercialstretching device (Zwick 1425). Stretching was per-formed with a rate of 4 mm/min and the appliedforce and elongation was recorded. Force was mea-sured with a 500 N load cell (accuracy: ±1 N). Forrecording the elongation, two optical markers wereattached on the middle of the sample at a distanceof 40 mm (investigated region). The lengthening ofthat region was measured by an optical detectionsystem (accuracy: ±0.5 mm). The cross section ofthe samples was determined with a micrometerscrew (accuracy: ±0.02 mm). Tensile curves are pre-sented as nominal stress, which is the measuredforce divided by the initial cross section as functionof strain (k = L/L0), where L is the actual and L0
the initial distance of the optical markers. First thesamples (n = 5) were stretched until failure to deter-mine the failure strain (kf) defined as the strain atbreak. P1, P2, P3 and P4 showed failure strainskf = 10.25; 9.7; 12.1 and 9.75, respectively. To inves-tigate the elastic behavior of the materials P2, P3and P4, strips were submitted to cyclic tensile tests.For the first cycle the samples were stretched to 80%of the previous determined failure strain (kf). After30 min of relaxation, the samples were stretchedagain in further cycles. Elasticity was determinedby the elastic recovery (reversible deformation/%)after the first stretching cycle by the formulaer = (kmax � krel/kmax � 1) · 100, where krel is therelaxed elongation after stretching to the maximalelongation kmax.
2.3. WAXS and SAXS
Wide-(WAXS) and Small Angle (SAXS) X-rayScattering measurements were performed at roomtemperature at the polymer beamline A2 at HASY-LAB at DESY in Hamburg. Phosphor imaging-plate devices (IP) were used for recording the twodimensional (2d) scattering intensities. For thestretching of the samples a custom-made setup,which is implemented in the beamline, was used.Samples were stretched stepwise from k = 1 tok = 7; 9.7 and 7 for P2; P3 and P4, respectively.At each step, a 2d scattering record was taken. Sam-ple elongation was determined using radiolucentmarkers on the strips and a caliper. Constant azi-muthally intensity distributions of 2d scattering
images assured, that only initially isotropic sampleswere investigated. For determination of the crystal-linity, a separation of crystalline and amorphousphase from the WAXS scattering was necessary[61]. Crystallinity was calculated from the WAXScurves after background correction. Data from theX-ray amorphous sample P4 were taken to deter-mine the amorphous scattering called herein asamorphous phase. The scattering from the crystal-line phase is than the absolute scattering minusthe amorphous. The crystallinity equals the integralover the crystalline scattering divided by the integralover the total scattering. To calculate the crystallin-ity, radial intensity distributions (h = 0–15�) wereused from the 2d records of isotropic samples(Fig. 1). 2d-WAXS records taken at different elon-gations are presented as colored pictures for visual-ization. For a quantitative description of theorientation, the radial intensity distributions wereread out (Fig. 1) and the orientation functions werecalculated. Radial read out starts from the positionof the primary beam along the stretching direction(Fig. 1, top down, u = 0�) to the equatorial direc-tion (u = 90�) in steps of 2� for each elongation.The width of the radial cross section was du = 1�.The scattering angle h was calculated using the
A. Boger et al. / European Polymer Journal 43 (2007) 3573–3586 3577
measured Bragg angles, setup dimensions, wave-length of radiation (1.5 A) and known formula[58]. The position of the primary beam was deter-mined by the symmetries of the picture. Amorphoushalo (scattering curve of the sample P4 in the isotro-pic state, Fig. 4) and background scattering (as linear
Fig. 2. Illustration of the transformation of the azimuthallyintensity distribution from 2d record (circle A) to the circle B onthe sphere for the calculation of the orientation function.
0.45
0.55
0.65
0.75
0.85
0.95
0 0.2 0.4 0.6
angle p
inte
nsi
ty /
a.u
.
L=1L=2L=3L=4.2L=5L=6L=7L=9.7
2θ = 13.65º
Fig. 3. Illustration of the extrapolation of the intensity distributi
curve) were subtracted from each radial intensity dis-tribution individually. Combining data of radialintensities at fixed scattering angles results in azi-muthally scattering intensity distributions. These azi-muthally intensity distributions were carried out foreach elongation caused by the crystalline structureof the material; I(hi = const., k, u), 1 6 k 6 kmax,0� 6 u 6 90�, 0� represents the stretching direction.Where the hi’s are the scattering angle of interest(Table 2), u the radial angle and k the elongation.The notations of the reflections given in Table 2 cor-respond to the identification of the apparent crystal-modifications [50].
Because the orientation function is defined on thesphere, the azimuthally intensity distributionsreceived from the planar detector have to be trans-formed to obtain the distribution on the sphere bytransforming the angle u/rad to the angle u 0/rad(Fig. 2). In order to calculate the orientation func-tion, a rotational symmetry of the texture, resultingin same intensities on each meridian was assumed.Therefore the intensity distribution on a meridianexcept for the region 0 6 u 0 6 h could be calculatedusing the azimuthally scattering intensity at fixedscattering angle h. The transformation formula isgiven by u 0 = arcos(cos(h) � cos(u)) [59], as illus-trated in Fig. 2. Intensities for 0 6 u 0 6 h wereadapted using regression (Fig. 3). Afterwards,the orientation function was calculated using the
0.8 1 1.2 1.4
hi`/ rad
on on the sphere (circle B, Fig. 1) for the range 0 6 u 0 6 h.
Fig. 4. WAXS scattering curves of samples in Table 1; k(X-ray) = 1.5 A.
Table 2Bragg angle (2h) and net plane identification of the interestpositions for calculation of the azimuthal intensity distributionfor the a–c modification according to [47]
Scattering anglesof interest in 2hi/�,k (X-ray) = 1.5 A
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formula below (Eq. 1, [60,61]) and the intensitydistributions I(h = const., k = const., u 0) = I(u 0).Related investigations on the orientation in PP filmswere done by Zigmond et al., where the method isdescribed in more detail [60,61]. Because the adapta-tion of the amorphous scattering to the scatteringcurve was performed visually, an error of 5% wasassumed. This assumption results as difference inorientation function values of several amorphousphases adjusted with the same subjectiveimpression.
f ¼ 1
2ð3 � hcos2ð/Þi � 1Þ; ð1Þ
where
hcos2ð/Þi ¼R p
20
Iðu0Þ cos2ðu0Þ sinðu0Þdu0R p
20
Iðu0Þ sinðu0Þdu0
For example the azimuthally intensity distributionfor the fixed scattering angle 2h = 18.2� correspond-ing to the (130)a net planes represents I(1 3 0)a(u 0),which was used for the calculation of the orientation
function f for the net planes (130)a. Orientationfunctions are presented as functions of elongation.The co-domain of the orientation function rangesfrom �0.5 (net planes are perpendicular orientedto the reference axis) to 1 (net planes are paralleloriented to the reference axis). Reference axis wasthe stretching direction indicated by u = 0�. Zero(f = 0) indicates an isotropic distribution of thenet planes. A rotational symmetry can be assumedif the sample dimensions are several times biggerthan oriented structures according to [62]. This isthe case for our polymers, where the sample thick-ness was greater than 0.1 mm even at maximumelongations. Obtained 2d-WAXS records were com-pared to the calculated ones from Auriemma andDe Rosa [51] for the preferred orientations ofa- and c-crystals including mother and daughterlamellae.
Additionally, the superstructures of the sampleswere investigated under uniaxial deformation per-forming SAXS measurements. The SAXS recordswere made similar to the WAXS records with longerdetector-sample distance and higher exposure time.SAXS images are presented as colored pictures forvisualization and discussed according to a modelwhich was described by Bratrich et al. [63]. Longperiod and block size according to the mentionedmodel were determined for directions obtaining dis-crete small-angle X-ray scattering.
3. Results and discussion
Table 1 lists the chemical and physical character-istics of the investigated PP samples, showing thewide range of used materials. Sample P1 was usedto give an impression of the mechanical propertiesof a higher crystalline, more conventional PP incomparison to the less crystalline, elastic PP samplesfocused herein.
3.1. Uniaxial stretching
As is apparent in Fig. 5, the sample P1 shows aclassical behavior for a high crystalline thermoplas-tic material in the tensile test, with a pronouncedyielding point and necking. The elastic recovery pre-sented by the strain start point for the 2nd loadingcycle reveals, that the elasticity increases with lowercrystallinity of the samples. The elastic recovery er
for the samples P1, P2, P3 and P4 was 5%, 37%,77% and 92%, respectively. Instead, the nearlyX-ray-amorphous sample (P4) achieves a complete
Fig. 5. Tensile tests of the investigated samples: P1 (upper left), P2 (upper right), P3 (lower left) and P4 (lower right). The grey curves arethe second cycle.
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resetting of the strain after the first drawing cycle(Fig. 5). The mechanical behavior of the sampleP4 is similar to that of a cross-linked rubber mate-rial [64], which can be described by the Van derWaals theory model for polymers by Kilian et al.[65]. Failure strain for P2, P3 and P4 was on averageef = 9.7, 10.25 and 9.75, respectively. Stress–straincurves of the hysteresis measurements of the sam-ples P2, P3 and P4 are reported in Fig. 5.
3.2. Two-dimensional WAXS
Fig. 6 shows WAXS-images of the sample P2 atvarious elongations from k = 1 to 7. Figs. 7 and 8show the WAXS-images taken during stretching ofthe samples P3 and P4. The WAXS measurementsshowed very distinctive fiber textures for all threesamples presented. Therefore, all samples investi-gated show crystal orientation during stretching ingeneral.
For sample P2 it could be qualitatively observed(Fig. 6), that at a deformation of k = 2.5, scatteringgets anisotropic due to orientation of the existingcrystals. Orientation functions calculated for the
sample P2 of the main net planes (110)a, (040)a,(130)a, (117)c and (202)c are shown in Fig. 9. Atk = 1 an actual value of nearly zero shows the ini-tially isotropic state of the sample. During stretch-ing, net planes are orienting towards the stretchingdirection (f < 0) and becomes more accurate athigher elongations, except the (040)a. Net planescorresponding to (117)c show a slight orientation(f(117)c = �0.08, k = 7) even at maximum elonga-tion. However, the orientation function of the netplanes (110)a and (130)a decreases to �0.33 atmaximum elongation indicating a medium orienta-tion along the stretching direction. Net planesdefined by (040)a initially show a slight orientationperpendicular to the stretching direction (f = 0.23,k = 2.5) and decrease afterwards to (f = �0.08,k = 7). The orientation of the (040)a net planes thatshould occur exclusively on the equator are notcompletely understood. Maybe these reflexes arecaused by tilted crystal structures (e.g., lath-likecrystals) or the reflexes are overlapped with oneanother at this scattering angle region ((1 10)a,which can be present in several orientation direc-tions in the aa-axis + ca-axis orientation [51]).
Fig. 6. WAXS images of the sample P2 at elongations k = 1, 1.25, 2.5, 4, 5, and 7 (in read succession); k (X-ray) = 1.5 A. Strain directionis top down in the pictures.
Fig. 7. WAXS images of the sample P3 at elongations k = 1, 4, 6, and 9.7 (in read succession) k (X-ray) = 1.5 A. Strain direction is topdown in the pictures.
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Fig. 8. WAXS pictures of the sample P4 at elongations of k = 1, 4, 5 and 7 (in read succession) k (X-ray) = 1.5 A. Strain direction is topdown in the pictures.
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Orientation of the (202)c net planes stays more orless in the isotropic state (f = 0) during stretchingthe material. The stress–strain curve of P2 (Fig. 5)shows an increase in strength with nearly constantslope after reaching the yield point. Compared tothe results from orientation investigations, this con-stant slope is associated with constant increasingorientation of main net planes ((1 10)a, (130)a) ofsample P2, which consists mainly of a-crystals[50]. Therefore, we conclude for P2, that the stiffen-ing effect comes from the increase of oriented crys-tals. Crystal orientations presented by P2 with itsrelative high crystallinity of 36% seems to be mainlyirreversible which can also be seen by its low elasticrecovery. Comparing the reflex intensities and loca-tions from the WAXS records to the model calcula-tions [51], we can conclude that main crystalsorientate during stretching according to a pure ca-axis orientation and to a lower content in a mixtureof aa-axis + ca-axis orientation.
Sample P3 indicates crystal orientation andrecrystallization at elongations lower than k = 4(Fig. 7). Quantitative calculations (Fig. 10) showedan increase in orientation for k > 4 of the main net
planes along the stretching direction to a low(f = �0.06), slight (f = �0.14) and medium(f = �0.31) extend for (110)a, (040)a and (202)cnet planes, respectively. As apparent in Fig. 10,most changes are between k = 3 and k = 5. Fig. 5indicates an increase in stiffness (slope of stressstrain curve) during the first stretching cycle espe-cially in the mentioned region and continues forhigher elongations. Comparing the slope in thestress–strain curve and the orientation it can be con-cluded, that the increase in cross-linking density(increasing slope/material stiffness) correlates withan increase in crystal orientation. This indicates,that these oriented crystals cause the revertive driv-ing force, which is the reason for the elasticity afterstress is released. Thus, they are the building blocksof the network responsible for material elasticity.The 2d WAXS images showed, that the equatorialreflexes ((008)c + (040)a, (111)c + (110)a) havingcomparable intensities, kept separated duringstretching with higher intensity as the third one((11 7)c + (130)a). The (202)c reflex appears strongat the first layer line. Therefore, sample P3 showed aparallel chain axis orientation of the disordered
Fig. 9. Orientation functions on elongation from the (110)a,(040)a, (130)a, (117)c and (202)c net planes of the sample P2.
Fig. 11. Orientation functions on elongation from the (110)a,(040)a and (202)c net planes of the sample P4.
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a/c-crystals according to the calculations fromAuriemma and De Rosa [51]. As apparent in sampleP2 and P3, the orientation function of the net planescalled (040)a showed a high correlation to the stiff-ness of the material and have to be analyzed andunderstood in more detail. The X-ray amorphoussample P4 shows pronounced strain-inducedchanges in the chain orientation/crystallization,demonstrated in Fig. 8 as it is known for conven-tional vulcanized rubber materials. The correlationbetween the orientation function of the net planesand the development of the slope in the stress–straincurve as already mentioned for the sample P3 can beobtained more pronounced for the sample P4
Fig. 10. Orientation functions on elongation from the (110)a,(040)a and (202)c net planes of the sample P3.
(Fig. 11). High correlated changes in the orientationand the increase in strength can be obtained at elon-gations at k > 4, which allow the same conclusionsas for sample P3. However, the absolute orientationof that material was lower due to the fewer amountsof crystals in the initial stage. Due to the broaderreflexes of sample P4 it is difficult to address theorientation patterns according to the mentioned cal-culations [51].
Additionally WAXS records were performed onhysteresis measurements. This investigation show,that if the stress is released, the crystals are stablebut relaxed in the orientation.
3.3. Two-dimensional SAXS
Discrete small-angle scattering could only beobtained for sample P2. Bratrich et al. [63] basedthe interpretation of the superstructure transitionson calculations of their anisotropic two-phasemodel.
The 2-dimensional SAXS records of the sampleP2 are shown for the elongations k = 1 to k = 6(Fig. 12). At k = 1 an isotropic distribution of thelamellae structure can be obtained. By stretching,the lamella bl1ocks are sliding to each other, result-ing in four-point SAXS diagram. The sliding canalready be seen at the first elongation step atk = 1.25 (Fig. 12). The four-point diagram shiftstowards the equator with increasing draw ratio. Athigher elongations the lamella blocks are separatedand a fibril structure is created. This can be seenat the meridian reflexion, which appears at k = 3.Up to an elongation of k = 6 the meridian reflexes
Fig. 12. SAXS pictures of the sample P2 at elongations k = 1, 1.5, 2, 3, 4 and 6. (k (X-ray) = 1.5 A).
A. Boger et al. / European Polymer Journal 43 (2007) 3573–3586 3583
become stronger, which signifies a distinct fibrilstructure. While stretching the sample to higherelongations, the sample gets white due to crazesand cracks.
The calculated long period in the isotropic SAXSrecord (Fig. 12, k = 1) was 122 A. Dimensions dueto structure transitions obtained in Fig. 12 can beexplained according to the model in Bratrich et al.[63]. The four-point diagram with its discrete scat-tering corresponds to regular structures with a per-iod of 133 A, 137 A and 180 A at k = 1.5, k = 2 and
k = 3, respectively. The lamella block size of fibrilstructures can be evaluated at higher elongationsaround k = 3 when they are separated. Block sizesof 180 A with a period of 128 A were determinedat k = 6 (Fig. 12).
SAXS investigations confirmed the well knownphenomena, that a certain amount of initial crystal-linity is required for building a superstructure. Acomparison of 2d SAXS records with sketches fromthe model shows a good conformity of the modeland the experimental results [63].
3584 A. Boger et al. / European Polymer Journal 43 (2007) 3573–3586
4. Summary and conclusion
The presented wide-angle measurements showedvery distinctive fiber textures. Thereby, the changein crystallization became apparent and, in particu-lar, a strain-induced crystallization for the low crys-talline sample (P4) could be observed. Using theazimuthally intensity distributions of the wide-angleimages, recrystallization and stress-induced crystal-lization could be determined. The orientations ofthe net planes were calculated by the orientationfunction, whereby a fiber texture was assumed.The orientation of the crystals, existing in the non-stretched state or induced by stretching, shows ahigh correlation with the stiffness in the stress straindiagrams. Higher stiffness values in elastic materialscorrespond to a higher density of cross-links.Answering the main question of the study aboutthe nature of the network in the high elastic PPmaterials (P3, P4) it is obvious that the crystaldomains in the soft amorphous matrix are responsi-ble for the elastic behavior and for building up thenetwork. Results of the present work support theinterpretation of De Rosa et al. [6] about the originof elasticity of low isotactic PPs. Orientation pat-terns observed were compared according to theworthwhile model calculations of Auriemma andDe Rosa [51] and where confirm with them.
Investigations of the superstructure in sample P2under uniaxial drawing revealed clearly the transi-tion from a lamella structure to a fibril structure.The interpretation of the superstructure transitionswas based on calculations of the anisotropic two-phase model.
The mechanical behavior of the samples coverselongation up to 20 times of the initial length anda recovery from 70% or rather 8 times with a100% (perfectly elastic) recovery after the first cycle.This shows the great potential of these PPs as ther-moplastic rubber materials with high recovery andsmall hysteresis.
The discussion of the orientation in terms of unitcells and stress strain measurements showed that thenetwork, which is responsible for the elasticity, ismainly composed of strain-induced crystallites.
Acknowledgement
The authors thank Stefan Fischer, UniversityUlm for proof-reading the manuscript and helpfuldiscussions.
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