Investigation of structural changes in semi-crystalline polymers during deformation by synchrotron...

Investigation of Structural Changes in Semi-Crystalline Polymers DuringDeformation by Synchrotron X-Ray Scattering

KONRAD SCHNEIDER

Leibniz Institute of Polymer Research Dresden, Hohe Straße 6, 01069 Dresden, Germany

Received 30 September 2009; revised 23 November 2009; accepted 23 November 2009

DOI: 10.1002/polb.21971

Published online in Wiley InterScience (www.interscience.wiley.com).

ABSTRACT: The mechanical behavior of polymer materials is

strongly dependent on polymer structure and morphology of

the material. The latter is determined mainly by processing

and thermal history. Temperature-dependent on-line X-ray

scattering during deformation enables the investigation of de-

formation processes, fatigue and failure of polymers. As an

example, investigations on polypropylene are presented. By

on-line X-ray scattering with synchrotron radiation, a time re-

solution in the order of seconds and a spatial resolution in the

order of microns can be achieved. The characterization of the

crystalline and amorphous phases as well as the study of cavi-

tation processes were performed by simultaneous SAXS and

WAXS. The results of scattering experiments are comple-

mented by DSC measurements and SEM investigations. VC 2010

Wiley Periodicals, Inc. J Polym Sci Part B: Polym Phys 48:

1574–1586, 2010

KEYWORDS: mechanical properties; on-line X-ray scattering;

polypropylene; SAXS; strain-induced crystallization; structural

characterization; WAXS

INTRODUCTION The stress-strain curves of semicrystallinematerials generally show three characteristic regions.Although the initial region before the yield point apparentlybehaves elastically, stress relaxation due to rearrangementsin the amorphous phase can be observed. The mobility inthe amorphous phase depends on the difference between theambient temperature and the temperature characteristic ofthe glass transition, which is the dominant relaxation processin the temperature range under investigation. After the yieldpoint, typically local necking occurs with high local strain,whereas the overall strain remains moderate. In the case ofdog-bone specimens then the neck propagates over thewhole specimen during constant load. In the true stress-strain diagram necking is a fast local deformation from theyield strain to the strain of the fully yielded specimen. In thethird step during further elongation strain hardening occurs,until finally the specimen fails.

Already in his first model of structural changes during defor-mation, Peterlin1 discussed spherulites consisting of lamellaeseparated by amorphous regions. Because of tie moleculesstress is transferred between the lamellae inducing deforma-tion and finally disintegration and rearrangement in theform of fibrillae. Later numerous authors modified andenhanced Peterlin’s approach to describe their experiments.The discussion gained new impetus after the first onlinestructure investigation during deformation by Chen et al.2

Breese and Beaucage3,4 gave an overview of older models.Finally, they describe mechanical behavior with a model

refining the models of Weeks and Porter5 and Gibson et al.6

in the sense that the effects of the orientation process on themodulus of the nonfibrous gel component are incorporatedinto the model by allowing a transition to fibers.

To describe the deformation behavior of semicrystalline poly-mers Strobl and coworkers7,8 separated different mecha-nisms of stress transfer with respect to amorphous and crys-talline units. They distinguished between four phases: Onsetof local block sliding (1), collective motion slightly below theyield point (2), disassociation of crystalline blocks and trans-formation into fibrils (3), and the start of disentangling (4).

In most models, the common process of cavitation in poly-mers during deformation is not yet incorporated. Many sys-tems show a characteristic whitening during plastic deforma-tion due to the formation of voids or cavities with a typicallength scale growing up to the wavelength of visible light,that is, some hundreds of nanometers (see Pawlak and Gale-ski9 and Men et al.10). It is obvious that these large scalestructures are not accessible with a conventional small angleX-ray scattering setup. But at least the beginning of this pro-cess should be well-detectable by X-ray scattering due to thestrong difference in electron density between the polymerand the voids in the relevant angular range accessible byUSAXS (see Lode et al.,11 Gehrke et al.,12 and Roth et al.13).

In previous studies, Davies et al.14 demonstrated a continu-ous generation of the voids in the plastic phase starting atthe yield point, which led to a conspicuous increase in the

Correspondence to: K. Schneider (E-mail: [email protected])

Journal of Polymer Science: Part B: Polymer Physics, Vol. 48, 1574–1586 (2010) VC 2010 Wiley Periodicals, Inc.

1574 INTERSCIENCE.WILEY.COM/JOURNAL/JPOLB

scattering power. Furthermore, a change in the diffuse scat-tering profile indicates a monotone change in the size and inthe shape of the voids.

Stribeck et al.15 described nanostructure evolution in poly-propylene during online mechanical testing with simultane-ous SAXS and refined details of the interaction between thedifferent phases including cavitation.

Processing conditions strongly affect the mechanical proper-ties caused by the interaction of the different phases withinthe samples. For instance in an injection molded plate, theproperties can be extremely dependent on the position anddirection of a specimen, changing from relatively brittle tohighly stretchable.

As an example the properties of isotactic polypropylene sam-ples, produced by compression and injection molding,respectively, were investigated under stepwise loading.Simultaneously, the structural changes were characterized bySynchrotron-SAXS and discussed together with the mechani-cal properties. It was found that there is a highly preorientedstructure within the injection molded specimen, whichstrongly influences the further temperature-dependent defor-mation behavior. The influence of temperature and preorien-tation on the deformation behavior is described extendingthe present models of plastic deformation.

EXPERIMENTAL

General ProcedureTo investigate structural modification during deformation X-ray scattering experiments were performed on samplesmounted in a miniaturized tensile rig placed in the synchro-tron X-ray beam. The synchrotron radiation is necessary tohave sufficient intensity to get high time resolution. A gen-eral description of the equipment was given by Davieset al.14

The actually used experimental arrangement and the speci-men geometry are illustrated in Figure 1.

To investigate local strain-dependent properties, smallwaisted specimens were used to concentrate the stress in

the center of the specimen. By using a relatively large radiusof curvature (12 mm) compared with the specimen width (3mm), the stress state is in good approximation maintainedconstant in the middle of the specimen. The strain wasdetermined optically by observing the deformation of a gridpattern applied on the specimen surface. The grid patternwas applied using a self-made flexible ink and a mesh sizeof 0.35 mm. The middle of the specimen, where the beamcrosses, was left blank. Alternatively, also a classical imagecorrelation analysis for strain estimation can be used. A com-parison of stress-strain curves of the waisted specimenswith standardized dog-bone specimens show a good consis-tency: The curve progression from the beginning to the yieldpoint as well as during strain hardening, that is, where theparallel region of the dog-bone specimen is deformed in auniform way, are identical. For the range of neck formationand propagation in the dog-bone specimen, an estimation oftrue stress and strain via global strain measurement are notpossible. Only a local strain measurement solves the prob-lem. The local measurement of the waisted specimens pro-vides true stress and strain data in a good approximation.

In this context, the specimens are assumed as incompressi-ble. This allows calculating the true stress rt from the meas-ured force F, initial cross section A, and tensile strain et as rt¼ F � (1þet)/A.

To characterize the structure at certain strains mostly step-loading experiments were performed, stretching the sampleto a certain strain and then recording the patterns at thisconstant strain. This is also the reason for the presentedslightly fluttering stress-strain curves. In some cases alsocontinuous stretching was performed.

To keep the beam in a fixed position relative to the gaugelength of the waisted sample throughout the measurement,both grips were moved simultaneously in opposite direction.

To get simultaneous SAXS and WAXS patterns, the WAXScould be monitored only in a limited range. By using a hori-zontal tensile direction, the vertically arranged WAXS detec-tor monitors mainly the equatorial scattering of the sample.

For the validity check, the sample was rotated around thetensile direction and the patterns were compared. To detect

FIGURE 1 Sketch of the experimental arrangement for SAXS and WAXS during deformation (left) and waisted specimen (mini-

dumbbell) used for simultaneous structure and mechanical investigation (right), the dimensions of the specimen can be scaled.

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X-RAY SCATTERING OF SEMI-CRYSTALLINE POLYMERS, SCHNEIDER 1575

oriented crystallites, the sample can be rotated around anaxis perpendicular to the beam.

To perform temperature-dependent tensile and scatteringexperiments, a small heating device locally blows preheatedair at the sample.

Besides the global scanning of samples, the experimentalsetup also permits spatially resolved pattern recording, forexample, around a crack tip, if it is used in a microfocusbeamline. The described arrangement was successfully usedinvestigating semicrystalline polymers during deformation(Schneider et al.16) and fracture (Schneider17,18).

WAXS and SAXS measurements were performed at the syn-chrotron beamline BW4 at HASYLAB in Hamburg, Germany.13

The wavelength of the X-ray beam was 1.3808 Å, the beamdiameter was about 400 lm. The SAXS images were col-lected by a two-dimensional MarCCD-detector (2048 � 2048pixels of 79.1 � 79.1 lm2). The sample-to-detector distancewas set to 4080 mm. The WAXS images were collected by atwo-dimensional PILATUS 100K-detector (487 � 195 pixelsof 172 � 172 lm2). By a special procedure, the position ofthe tilted WAXS-detector was determined to have a tilt angleof 22.7�. The distance between the sample and the point ofnormal incidence was 247 mm.

Exposure times were chosen in the range of 5–60 s per pat-tern. The frame rate, determined by exposure and data stor-age, was 15–70 s per pattern.

For the discussion, all SAXS and WAXS 2d-patterns areshown with vertical tensile direction. For specimens withfiber symmetry, this means that the fiber axis and so thescattering vector s3 also are vertical, the scattering vector s12always are horizontal.

For quantitative comparison of crystallinity, some DSC meas-urements were performed on drawn and undrawn samplesusing a DSC Q 1000 from TA-instruments. The specimenswere measured between �80 �C and 230 �C with heatingand cooling rates of 20 K/min. Before each scan, the samplewas equilibrated at constant temperature for 300 s.

As qualitative check of the discussed structures, some SEMimages of the drawn samples were taken on a Zeiss GeminiUltra plus SEM.

MaterialThe investigations were done on isotactic polypropylene(iPP), grade HD 120 MO from Borealis. Both injection andcompression molded plates were used. Melt temperature ininjection molding was 245 �C, mold temperature 40 �C. Themold has a film gate with a width of 0.5 mm, the size of theplate is 80 mm � 80 mm. The compression molded plateswere produced by heating injection molded plates to 190 �Cfor 5 min in a vacuum press. Then the plates were cooleddown to room temperature at 15 K/min. All plates have athickness of about 1 mm. Mini-dumbbell specimens wereproduced by CNC milling. From the injection molded plate,specimens were taken near the inlet in injection directionand perpendicular to the injection direction.

Evaluation of Scattering DataSAXS Data EvaluationThe evaluation of SAXS images is strongly related to theapproach for materials with fiber symmetry developed byStribeck.19 Data processing was realized with the softwarepackage pv-wave from Visual Numerics.

The images were normalized with respect to the incidentflux and blind areas were masked. Considering the sampleabsorption, the instrument background was subtracted. Bytranslation and rotation, the images were aligned in that waythat the tensile direction is vertical and the beam position inthe middle of the patterns. Finally, the patterns were aver-aged with respect to the four quadrants of the detector. Theblind area of the beam stop is interpolated assuming a Guin-ier-type behavior of intensity in this range.

Generally, SAXS realizes a projection of the specimen struc-ture into the reciprocal space, where it produces a slice ofthe Fourier transform (FT) of the structure whose amplitudeis measured. A single back-transformation would deliver aprojection of the autocorrelation function of the structure.

Complete information about the specimen would be avail-able only by tomographic methods with a stepwise rotationof the sample (see e.g., Schroer et al.20) or using inherentsymmetry properties of the sample. Under the assumptionof fiber symmetry of the stretched specimen around thetensile axis, from the slices through the squared FT-struc-ture, the three-dimensional squared FT-structure in recipro-cal space can be reconstructed and hence also the projec-tion of the squared FT-structure in reciprocal space. TheFourier back-transformation of the latter delivers slicesthrough the autocorrelation function of the initial structure.Stribeck pointed out that the chord distribution function(CDF) as Laplace transform of the autocorrelation functioncan be computed from the scattering intensity I(s) simplyby multiplying I(s) by the factor L(s) ¼ �4p2(s212 þ s23)before the Fourier back-transformation. Here s is the scat-tering vector with the component s12 in transversal direc-tion and s3 in fiber direction.

The interpretation of the CDF is straightforward,21 because ithas been defined by the Laplacian of Vonk’s multidimen-sional correlation function.22 It presents autocorrelations ofsurfaces from the scattering entities in that way that positivevalues characterize distances between surfaces of oppositedirection, negative values distances between surfaces of thesame direction, see Figure 2.

Hence, positive peaks in the vicinity of the origin character-ize size distributions of the primary domains. In the case ofsemicrystalline materials, this can be crystallites as well asamorphous regions in-between. If cavitation occurs, it will besuperimposed by the size of the cavities. Also a relativelysmall amount of cavities will be visible because the differ-ence in electron density is here much higher than betweencrystalline and amorphous phase within the polymer.

Following negative peaks characterize distances between ad-jacent repetition units (long periods).

JOURNAL OF POLYMER SCIENCE: PART B: POLYMER PHYSICS DOI 10.1002/POLB


Positive peaks at greater distances describe size and orienta-tion of domains (from the beginning of the first domain tothe end of the second one). In-well oriented systems alsopeaks of higher order can be observed.

For detailed discussion, the CDF’s can be presented as con-tour plots or as density plots in the plane.

WAXS Data HandlingAccording to the experimental arrangement, the WAXS dataunfortunately are restricted to only a certain range of angles.The position of the WAXS detector was estimated using cer-tain reference reflexes of iPP. The procedure developed forthis purpose will be published separately.

In a first step, the individual images were normalized withrespect to the incident flux and blind areas were masked.The instrument background was subtracted considering sam-ple absorption. The images were always transformed to ref-erence coordinate systems with respect to the scatteringangles (scattering angle 2H and azimuthal angle u or scat-tering vectors s12 and s3).

For further discussion, the scattering intensity was projectedover a selected range of u on the 2H-axis or discussed overa selected range of 2H on u-axis.

RESULTS

Preliminary Mechanical TestsBecause of the waisted sample geometry, the whole defor-mation of the samples is different compared with standar-dized dog-bone specimens. We no longer observe the for-mation of a neck which moves at constant load across theparallel region of the specimen. Once the neck has formed

the sample immediately continues stretching within thenecked region due to the increasing cross section of thesample adjacent to this region. The comparison of technicaland true stress-strain curves for the three specimens underconsideration is shown in Figure 3. The strain is measuredoptically. True stress is calculated assuming constant vol-ume as rt ¼ F � (1þet)/A, naturally it is always higherthan the technical stress.

The injection molded specimens in injection direction aswell as the compression molded and quickly cooled speci-mens are highly stretchable. By contrast, the specimenstransversal to the injection direction fail very soon. Here insome cases, even crazing could be observed before failure.This points out that there will be a strong structural anisot-ropy due to the processing history (shear stress as well ascooling rate).

Check of Fiber SymmetryThe evaluation of SAXS patterns uses the constraint of fibersymmetry. To check the validity of this constraint severalsamples were investigated under three different angles ofincident beam. By rotating the sample, the X-ray beam pene-trated the sample from the flat side, the narrow side, and45� in-between. Comparison of the normalized patternsshowed no significant differences, which would be expected,for example, in the case of any affine deformation in one ofthese directions. Small differences arise due to a skin-core-structure of the samples. With increasing strain, the differen-ces become smaller.

Crystallite IdentificationAccording to Bragg’s law, the positions of WAXS reflexesrefer to the distance between crystalline planes within thecrystallites. In the compression molded as well as in theinjection molded plates, a couple of crystalline reflexes couldbe identified. The reflexes within the relevant angular regionare summarized in Table 1.

FIGURE 2 Information of CDF’s: In any particular structure,

positive values of the CDF represent correlations of interfaces

with opposite direction (light arrows); negative values repre-

sent correlation of interfaces with the same direction (dark

arrows).

FIGURE 3 Stress-strain curves of waisted specimens from com-

pression and injection molded iPP specimen at room tempera-

ture. The technical stress is always the curve with the lower

stress values.

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Elastic Crystallite DeformationLoading of the samples below the yield point causes a cer-tain shift of the crystalline peaks. According to Bragg’slaw, the peak positions refer to spacing between latticeplanes. Their shift inversely reflects the strain vertical to thecorresponding plane. In this context, the shift of theobserved (h k l)-reflexes in equatorial direction (perpendicu-lar to tensile direction) display the transversal crystallinestrain of the crystallites aligned in tensile direction.

Cyclic loading and unloading gives an estimation of the elas-tic crystalline strain also during plastic deformation. Figure 4shows transversal crystalline strain of the (1 1 0)-reflex ver-sus the overall tensile strain for three consecutive loadingcycles. The first strain amplitude was 0.12, the remainingstrain after unloading 0.04. The following cycles wereconducted within the overall elastic region. The shift of other(h k l)-reflexes yield the same result.

In a range of up to about 8% strain, the crystallites deformtransversally by about �0.8%. The mismatch between crys-talline and global strain indicates the different stiffness ofthe crystallites compared with the whole specimen due tothe relatively soft amorphous regions. During further loading,the strain of the crystallites remains constant. This meansthat further deformation is realized on the microscale onlyby the amorphous phase and by additional shear dislocationof the crystallites.

Crystallite Reorientation During DrawingThe stress-strain diagrams captured during the scatteringexperiments are shown in Figure 5. The optical measuredstrain did not exhibit the stability as in the preliminary me-chanical tests mainly due to the stepwise loading, but thegeneral trend is representative. As expected, the stress-levelof the stress-strain decreases with increasing temperature.Otherwise, it is generally lower than in the preliminary testsdue to the generally lower strain rate.

During deformation, the WAXS patterns change in a quitecharacteristic way, see Figure 6. Initially the intensity of the

crystalline peaks is relatively constant over the azimuthalangle pointing to a homogeneous distribution of crystallitedirections. At deformations below the yield point, slightchanges in the intensity of the peaks are reversible.

Above the yield point, the deformation behavior is stronglytemperature-dependent. Generally in the investigated equato-rial direction only the (h k 0)-peaks remain and the scatter-ing intensity concentrates azimuthally in equatorial direction.This is generally an indication of orientation of crystalliteswith the c-axis (chain direction) in tensile direction.Although the peaks sharpen with increasing temperature,which is an indication of growing crystallites, they decreasestrongly and broaden at temperatures in the range immedi-ately above room temperature. This indicates a gradual dis-ruption of crystallites into smaller fragments. Finally, atroom temperature, the distinct peaks disappear but give abroad halo in the angular range of the previous peaks. Thisindicates a rough orientation of the chains or very smallcrystallite fragments in tensile direction, probably in theform of fibrils. However, the thermal mobility of the chainsseems to be all in all insufficient to establish newcrystallites.

TABLE 1 Crystalline Reflexes of iPPa

Reflex

Intensity

(Qualitatively) d/A

Scattering

Angle 2H

1 1 0 Strong 6.269 12.64

0 4 0 Strong 5.240 15.13

1 3 0 Strong 4.783 16.59

1 1 1 Strong 4.170 19.05

0 4 1 Very strong 4.058 19.58

1 3 1 Very weak 3.634 21.89

0 6 0 Weak 3.493 22.79

2 0 0 Very weak 3.281 24.28

2 2 0 Weak 3.131 25.46

a Crystalline reflexes were used for calibration of the position of the

WAXS-detector as well as for further discussion of changes during

deformation.

FIGURE 4 Transversal crystalline strain versus optically mea-

sured tensile strain of semicrystalline iPP specimen estimated

via the shift of the (1 1 0)-reflex at room temperature.

FIGURE 5 Temperature-dependent stress-strain diagrams of

the compression molded iPP specimens under investigation.



Transmission and Total/Maximal Scattering IntensityFor the flat samples under investigation, the transmissionwas 0.83 in the unstretched state. During stretching and si-multaneous thinning, the transmission generally increased to0.96.

SAXS During StretchingFrom the set of successive scattering patterns during the de-formation for the following discussion, only patterns and thecorresponding CDF’s were chosen where characteristicchanges could be seen.

The respective results of SAXS from stretching at room tem-perature are shown in Figure 7. The regular circular form ofthe CDF from the beginning is nearly unchanged to about6% strain also during unloading. It represents the randomlydistributed crystallites, which were deformed elastically, asdiscussed in the section about WAXS above. The diameter ofthe inner ring in the CDF indicates a lamellar thickness of9.8 nm. The isotropic ring in the negative direction repre-sents the long period of 17.7 nm. The following second posi-tive and negative reflexes indicate a certain correlation withthe third neighbors. Missing higher-order reflexes indicatethat there are surprisingly no remarkable correlations abovethe third neighbors. During deformation below the yieldpoint (loading, unloading, relaxation, and repeated loading),this general situation is unchanged.

Beyond the yield point, the situation changes drastically.First of all it is noticeable that the total scattering intensityas well as the extrapolated intensity I0 increase dramaticallysuggesting the activity of new strong scatterers—presumablycavities, which exhibit a strong scattering contrast to iPP due

to their extremely low electron density. They start as smallcracks transversal to the tensile direction and finally deformto long drawn cavities between fibrils. According to theirsuccessive growth and hence their broadly distributeddimensions, they hardly show distinct scattering signals, butoverlap with the signals arising from the lamellae or theirfragments. The growing cavities are also responsible for thewhitening of the cold drawn samples, see also Pawlak andGaleski.9 In contrast to the scanning experiments with amicrobeam reported by Roth et al.,23 which allows to resolvecertain individual voids, here over a large ensemble of cav-ities is averaged. Within the CDF’s, these cavities arereflected by the broad initial decay in fiber direction, athigher strains perpendicular to fiber direction, see Figure 7,CDF’s in log scale on the right.

Yet at a strain of 12.5%, the beginning of the yielding region,the structure becomes anisotropic. Although thickness of thelamellae remains constant, they align perpendicular to thetensile direction. The correlation to the next neighbors intransversal direction is lost and concentrated within a4-point-pattern. This indicates occasionally an internal sheardeformation. Under the external stress, the lamellae aresheared and finally break down into certain blocks asdescribed by Strobl and coworkers.7

A new transversal correlation length is established at about37% strain. By shearing, the lamellar blocks break down toa size which later represents the dimensions of the fibrils.According to the relatively low internal mobility, the alignedchains are not able to form new crystallites. Their meandimension in transversal direction is with 8.4 nm clearlybelow the lamellar thickness. This correlates with the WAXS

FIGURE 6 Equatorial scattering intensity of undeformed iPP at room temperature and during strain hardening at elevated tempera-

tures. The curves always show a projection of the angular range of 62� around the equatorial direction. Right: WAXS-patterns, azi-

muthal angle versus scattering angle 2H, at three characteristic points: Undeformed as well as immediately before failure at room

temperature and 150 �C. [Color figure can be viewed in the online issue, which is available at www.interscience.wiley.com.]

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results, which also indicate the final disappearance of the la-mellar transversal order during cold drawing. Ultimately, athigh strains, this transversal correlation dominates the wholeCDF.

The general changes within the oriented chains—withincrystallites as well as noncrystallized stretched chains—areshown in the sketch in Figure 8, deformation at room tem-perature is shown in steps (a–c), (f), and g).

With increasing temperature, the deformation behaviortotally changes. The patterns and the CDF’s of the stretchingof iPP at 130 �C are shown in Figure 9. Here the averagedand the extrapolated maximum intensity I0 does not changedramatically above the yield point. Instead soon the lamellaealign in tensile direction increasing their correlation in thisdirection, indicated by higher order of reflexes in the CDF’s.The transversal displacement of the reflexes suggests thatthe aligned lamellae are shifted against each other. The situa-tion does not change generally also at high strains. Thisinterpretation is also supported by the WAXS results

described earlier. The deformation is shown schematically inFigure 8(a–e).

With increasing temperature, also the whitening is remark-ably reduced indicating that due to the higher mobility ofthe amorphous phase nearly cavitation is negligible as com-pared with the extent observed at room temperature.

SAXS of Injection Molded SpecimenIn comparison with the compression molded specimens alsoinjection molded specimens were investigated. Samples werecut longitudinally and transversally to the injection directionnear the inlet of the plates of 1 mm thickness. The SAXSpatterns and the CDF’s during deformation are shown inFigures 10 and 11.

It is apparent that the CDF’s of the initial sample are compa-rable with the compression molded one stretched at 130 �C.The only difference is that there are nearly no small blocks,but larger lamellae aligned vertically to the tensile direction.

FIGURE 7 Deformation of iPP at room temperature, from left to right: SAXS patterns, positive and negative CDF’s (always log-

scale, pseudo color) as well as a surface plot of the CDF’s (linear scale and for two strains also log scale) at different strains e ¼0.0, 0.125, 0.37, and 5.3 (from top). Each square of the pattern covers a range �0.12 nm�1 < s1, s3 < 0.12 nm�1, each square of

the 2d-CDF covers a range of �100 nm < r12, r3 < 100 nm, fiber direction always vertical. The stretching (fiber) direction is indi-

cated in the surface plots by arrows.



FIGURE 8 Sketch of the transformations of the oriented chains during deformation, in-between are additional amorphous chains:

(a) Lamellae with some tie-molecules, (b) elastic shear-deformation of lamellae under small load and reorientation with respect to

the load, (c) fracture of lamellae into smaller blocks due to local stress concentration caused by tie molecules, (d) stretched

aligned, but not recrystallised chains between the bocks during stretching at higher temperatures, (e) some of the fibrillar

arranged molecules crystallize, final stage in the case of hot stretched PP, (f) further dissolution of the blocks creating more

extended chains at room temperature, and (g) finally, there are several strands of extended chains, not crystallized, with some

amorphous regions in between, final stage in the case of cold stretched PP. [Color figure can be viewed in the online issue, which

is available at www.interscience.wiley.com.]

FIGURE 9 Deformation of iPP at at 130 �C, from left to right: SAXS patterns, positive and negative CDF’s (always log-scale, pseudo

color) as well as a surface plot of the CDF’s (linear scale) at different strains e ¼ 0.0, 0.08, and 1.3 (from top). Each square of the

pattern covers a range �0.12 nm�1 < s1, s3 < 0.12 nm�1, each square of the 2d-CDF covers a range of �100 nm < r12, r3 < 100

nm, fiber direction always vertical. The stretching (fiber) direction is indicated in the surface plot by arrows. [Color figure can be

viewed in the online issue, which is available at www.interscience.wiley.com.]

Here a strong correlation in tensile (that means injection)direction is apparent. Only the characteristic length of thelamellae with 7.24 nm and a long period of 14.1 nm areslightly smaller than in the compression molded specimens.The later evolution of the patterns of the sample in injectiondirection correlates with the patterns of the cold drawn com-pression molded sample described earlier, however, the lamel-lar thickness slightly increases to 14.7 nm. The final fibrilwidth is 5.5 nm.

The sample transversal to injection direction changesdirectly from the transversal correlation caused by the pro-cessing induced orientation to the general fibrillar correla-tion pattern typical for the stretched compression moldedsample, see Figure 7, bottom row. Here the characteristicdimension of the fibrils is about 8.1 nm.

The WAXS patterns of the injection molded specimens showrings representing a random lamellae distribution and some

intensity maxima along the rings representing a preferentialorientation of some of the lamellae.

Over the whole deformation range stress relaxation isobserved as soon as the tensile rig was stopped. However,because this can be ascribed to the amorphous phase, SAXSresults remain mostly unaffected.

DSC Investigation of Unstretched andStretched SpecimensThe heating peaks of the first heating curves of the stretchedsamples are very different, see Figure 12. It is possible tosplit each curve into two individual peaks whose positionsshift with the stretching temperature to higher values. Fur-thermore, although the high-temperature peak is nearly con-stant, the low-temperature peak increases with stretchingtemperature.

FIGURE 10 Deformation of injection molded iPP in injection direction at room temperature, from left to right: SAXS patterns,

positive and negative CDF’s (always log-scale, pseudo color) as well as a surface plot of the CDF’s (linear scale) at different

strains e ¼ 0.0, 0.1, 0.81, and 2.81 (from top). Each square of the pattern covers a range �0.12 nm�1 < s1, s3 < 0.12 nm�1, each

square of the 2d-CDF covers a range of �100 nm < r12, r3 < 100 nm, fiber direction always vertical. The injection and stretching

(fiber) directions are indicated in the surface plot by arrows. [Color figure can be viewed in the online issue, which is available

at www.interscience.wiley.com.]



The two peaks indicate that due to stretching at differenttemperatures, certain different crystallites evolve, lamellarand fibrillar, with different extension. This is in accordancewith the scattering results. The lower melting peak shouldbe that of the fibrils. The frozen stresses in the samplesstretched at lower temperatures may shift the peaks addi-tionally to lower temperatures.

These curves are very sensitive to the individual piece of thestretched sample. This is also the reason for the outlier ofthe curve of the 130 �C-stretched sample.

The second heating curves of all stretched specimens arealmost identical, but they differ from the second heatingcurve of the unstretched sample, see Figure 13. Melting ofthe latter is finished almost 5 �C later. This means that dur-ing heating to 230 �C and annealing there for 5 min, thecrystallites indeed melt, but the global alignment of thechains remained. Therefore, again a fibrillated crystallinemorphology will be obtained by spontaneous crystallizationduring cooling.

When comparing the second heating of the samples with thefirst one shows that the stretched samples have up to 15%higher melting enthalpy, respectively, crystallinity, increasingwith the stretching temperature.

FIGURE 11 Deformation of injection molded iPP transversal to injection direction at room temperature, from left to right: SAXS pat-

terns, positive and negative CDF’s (always log-scale, pseudo color) as well as a surface plot of the CDF’s (linear scale) at different

strains e ¼ 0 and 0.59 (from top). Each square of the pattern covers a range �0.12 nm�1 < s1, s3 < 0.12 nm�1, each square of the

2d-CDF covers a range of �100 nm < r12, r3 < 100 nm, fiber direction always vertical. The stretching (fiber) direction is indicated in

the surface plot by arrows. [Color figure can be viewed in the online issue, which is available at www.interscience.wiley.com.]

FIGURE 12 DSC-measurements: First heating curves of com-

pression molded iPP-specimens stretched at different tempera-

tures. For clarification the curves are vertically shifted.

FIGURE 13 DSC-measurements: Second heating curves of

stretched and unstretched compression molded iPP-specimens.

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SEM of Stretched SpecimensSEM images were taken to verify the described general de-formation behavior. Figures 14–17 show the compressionmolded unstretched reference specimens and specimensstretched at room temperature, 90 �C and 150 �C. The speci-mens were cut in tensile direction, and vapor deposited withplatinum. Thus, stretching direction and s3-axis are alwaysvertical in the figures, the s1-axis horizontal.

Figures 14–17 show the voids in the stretched samples, theformation of which was discussed earlier. Their sizedecreases with increasing stretching temperature. In thesample stretched at 150 �C, they are not found. The lamellaeand fibrillae are not visible at the present magnification, onlya slight texture in stretching direction can be seen. A similartrend was reported recently by Pawlak and Galeski,9 whofound voids only at stretching temperatures below 60 �C, athigher temperatures only deformed shperulites and fibrils.

CONCLUSIONS

The experiments revealed the continuously changing mecha-nisms of elastic and plastic deformation and energy dissipa-tion (the transformations of amorphous and crystallinephases and void formation). Because of the high number ofparallel processes, the use of structure characterization by X-ray scattering techniques was realized in wide as well assmall angle range and complemented by DSC and by electronmicroscopy. This combination of methods is likely to providea well-founded basis for material development and optimiza-tion. DSC is an essential extension to characterize crystalliza-tion behavior.

It was found that there is a strong temperature-dependentinteraction of deformations within the amorphous phase andreorientations as well as transformations of crystalline units.They are strongly determined by the molecular structureand the processing dependent initial morphology of the sam-ples. According to the temperature-dependent mobilitywithin the amorphous phase, the stress transfer to the crys-talline phase and so the changes within this phase are alsostrongly temperature-dependent.

The mechanisms found and discussed here in the case of iso-tactic polypropylene seem to be very universal. It might alsoapply to other polymers. This will be checked in the futureby additional investigations.

The detailed study of the multiple deformation processes isan essential base for the development of semicrystallinematerials with well-defined properties.

This article is dedicated to M. Stamm in the occasion of his 60thbirthday. He had encouraged the author and his group to estab-lish the presented experimental and evaluation environmentand to perform the experiments.

Furthermore, the author is grateful to HASYLAB for beamtimewithin the project II-20060086, A. Timmann and St. Roth forlocal support, and N. Stribeck (University of Hamburg) for the

FIGURE 14 Scanning electron micrograph of the unstretched

reference iPP sample.

FIGURE 15 Scanning electron micrograph of the iPP sample stretched at room temperature. Stretching direction vertical.



support in the course of data evaluation. Finally, the authorthanksW. Jenschke for the software support for the tensile rig, A.Schone for support in data evaluation, R. Sommerschuh and D.Krause for sample preparation and preliminary tests, S. Frenzelfor the DSCmeasurements, R. Boldt for SEM images, and K. Brun-ing for revising themanuscript.

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FIGURE 16 Scanning electron micrograph of the iPP sample stretched at 90 �C. Stretching direction vertical.

FIGURE 17 Scanning electron micrograph of the iPP sample stretched at 150 �C. Stretching direction vertical.


ARTICLE

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