domain: spi phys

text:
J Martin

M Ponçot

P Bourson

A Dahoun

J M Hiver

Study of the crystalline phase orientation in uniaxially stretched polypropylene by Raman spectroscopy: validation and use of a time-resolved measurement method

Keywords: Isotactic Polypropylene, Uniaxial Tension, Deformation Mechanisms, Molecular Orientation, Relaxation, Wide Angle X-ray Scattering (WAXS), Raman Spectroscopy, Realtime measurements

A method based on the intensity sensitivity of two Raman scattering bands (973 and 998 cm -1 ) to laser light polarization is proposed to quantify the molecular orientation of the crystalline phase in isotactic polypropylene. This paper is firstly devoted to the validation of the method by comparing Raman and Wide Angle X-ray Scattering (WAXS) measurements performed on post mortem test samples uniaxially deformed at growing levels. The method is secondly applied to follow the molecular orientation in real-time during the course of a tensile test. Post mortem results show a quantitative agreement between Raman and WAXS measurements. In situ results mainly evidence the ability of the Raman method to characterize molecular orientation in real-time with less than 5 second time resolution. Comparison between post mortem and in situ results gives insight of relaxation mechanisms occurring in material after unloading.

INTRODUCTION

To gain a better understanding of the macroscopic mechanical properties of materials, especially for solid polymers, it is fundamental to understand the mechanisms of deformation which take place at the microstructure scale [1]. In the specific case of semi-crystalline polymers which exhibit a complex morphology mixing a disordered amorphous and an ordered crystalline phase, mechanical properties are mainly dependent on both the degree of crystallinity and the macromolecular orientation [1][2][3]. At the macroscopic scale, mechanical properties of such materials are usually determined using conventional uniaxial tensile tests [5]. At the microscopic scale, a number of analytical techniques can be used to investigate the microstructure such as thermal analysis, microscopic observations, X-ray diffraction, birefringence, nuclear magnetic resonance and vibrational spectroscopic techniques [6,7]. By collecting both these macroscopic and microscopic data, numerous experimental works have attempted to associate one state of deformation (such as the strain level) to one state of microstructure (such as molecular orientation) in order to propose elementary descriptive models of the deformation micro-mechanisms [8][9][10]. However, in many cases, evidence of these mechanisms is debatable for two major experimental limits: i) In the case of ductile materials which experience a pronounced necking due to their plastic instability (such as polymers), conventional mechanical tests giving access to the engineering stressstrain curve are inappropriate to reveal their intrinsic mechanical behaviour at prescribed temperature and strain rate.

ii) Moreover, usual analytical techniques require in most of the cases post mortem analyses and destructive sampling procedures coupled with serious risk of uncontrolled relaxational changes. In a recent paper, Tanaka and Young [] point out advantages and disadvantages of several techniques such as wide angle X-ray scattering WAXS, birefringence measurements, nuclear magnetic resonance and vibrational techniques for the characterization of the molecular orientation in polymers. Birefringence can be employed for rapid measurements but the orientation information is averaged over the crystalline and amorphous region. In the case of highly crystalline polymers, the measurement is also likely to be influenced by sample opacity. WAXS is a useful method for the study of crystallite orientation with accuracy. Using synchrotron radiations, this technique allows performing time-resolved measurements with less than 100 ms time resolution. Nevertheless access to beamline is often too restrictive for ordinary measurements. Focusing on the vibrational spectroscopic approaches, both infrared (IR) dichroism measurement and polarized Raman spectroscopy was found to be powerful techniques for quantitative and in real-time measurement of molecular orientation distribution both in amorphous and crystalline phase.

However, the major limitation of IR method is that samples have to be highly transparent and thin for the absorbance measurement.

In this work we propose to overcome these two experimental difficulties by combining both a tensile system which provides the true mechanical behaviour of a material and a polarized Raman micro-spectrometer. This systeme is used to characterize the molecular orientation in crystalline phase of isotactic polypropylene. The study is divided into two parts.

The first one deals with the validation of a specific measurement method based on the intensity sensitivity of two specific polarized Raman bands (at 973 and 998 cm -1 ) to orientation of crystalline chains. Method has been qualitatively used in a previous paper to discriminate orientation of lamellae within one -iPP spherulite. Its reliability is examined by comparing Raman results with the well-established Hermans' orientation function calculated from WAXS measurements performed on the same post mortem stretched samples. The second part of the study is attached to demonstrate the ability of the method to measure orientation in situ with the deformation (5 s time resolution) over a large true strain range, from the elastic to the hardening stage of iPP.

EXPERIMENTAL DETAILS

Material

The investigated polypropylene was manufactured by Atochem under the reference 3050 MN1. Molecular weight distribution of this isotactic grade was investigated by gel permeation chromatography. Polypropylene was dissolved in trichlorobenzene at 150 °C in order to obtain a 2 g.l -1 concentrated solution. Solution was then stabilized with ionol.

Analyses were performed at 145 °C. IPP is characterized by a relatively broad molecular weight distribution with M w = 262000 g.mol -1 and M n = 75900 g.mol -1 . Pellets of polymer were processed by injection-moulding into rectangular plates (2002004 mm) as shown in figure 1. The process consists of filling the mould maintained at 30 °C with the polymer melt at 240 °C, under a pressure of 1200 bars. The cooling sequence is about 240 s. As it is well reported in literature, injection-moulding process induces the development of an oriented superficial layer, usually termed as a skin layer []. It is the result of a shear stress at the wall of the cooled mould. Pictures of the figure 2, recorded with a polarized light optical microscope on a thin cross-sectional slice (20 µm thickness) extracted from the center of a plate evidences an anisotropic superficial layer of about 360 µm thickness. The core region is isotropic with an important amount of -spherulites of about 55 µm average diameter. In order to prevent measurement interferences with the molecular orientation induced by the deformation, skin layers of each specimen are previously eliminated by polishing flat surfaces. Polishing procedure is maintained identical for each specimen. Moreover to insure reproducibility, specimens are always extracted from the same location in the plate namely in its center far-away from the injection gate and the walls of the mould (figure 1). Material density in the core region is determined by hydrostatic weighting,  = 0.908 g.cm -3 . Based on values published in the literature for perfectly crystalline and amorphous phase ( c = 0.946 g.cm -3 and  a = 0.854 g.cm -3 , respectively [24]), the weight degree of crystallinity is found to be equal to 59 %. This value is confirmed by differential scanning calorimetry measurements performed at a heating rate of 10 °C.min -1 . Thermograms show a melting point located at 165 °C with an enthalpy of fusion equal to 96 J.g -1 . This corresponds to a weight degree of crystallinity of about 58 % (the enthalpy of fusion considered for a perfect crystal of infinite size is 165 J.g -1 [24]). X-ray scattering measurement performed in the core region with a Cu-K 1 radiation ( = 0.1542 nm) proves the semi-crystalline nature of the investigated iPP.

Narrow peaks observed at 2 = 14.1, 16.9, 18.5, 21.2 and 21.8 degrees show the presence of -iPP crystals in majority. They are due to the selective diffraction of the (110), (040), ( 130), ( 111) and ( 31) crystallographic planes of the -iPP monoclinic unit cell respectively [25].

The broad halo noted at 2  16.5 degrees is attributed to the X-ray diffusion of the amorphous phase.

Test specimens are machined out from the injection-molded plates into dumbbellshape specimens with overall dimensions of 90 x 10 x 4 mm 3 (Figure 1). As recommended in previous works using the VideoTraction TM system, the median region is reduced to 9 mm in order to localize necking in this predefined zone, the Representative Volume Element (RVE) [26,27]. Massive heads on both sides are designed to be gripped in jaws of the tensile machine.

Post mortem measurements

The main objective of post mortem measurements is to define the Raman measurement method and validate its use. For this purpose, specimens are stretched and unloaded at various strain level inducing different degree of the molecular orientation. Last one is firstly determined by WAXS measurements and then compared to Raman results.

The video-controlled mechanical testing method used here is the VideoTraction  which was developed by G'Sell and Hiver [29]. Figure 3 is a schematic view of the system. It gives access to the constitutive plastic behaviour of polymers under uniaxial tension by measuring in real time the true strain  zz (even after the necking phenomenon has appeared) and the true stress  zz in a given local material element (Representative Volume Element RVE or necking zone) while its local true strain rate is maintained at a constant value during the test. Technical explanations of the method principle were widely presented elsewhere before [14,15]. In our study, all tests are carried out at 30 °C under a monotonous strain rate maintained at zz ε . = 5.10 -3 s -1 . Various specific true strain values are chosen. They are summarized in table 1. Under load they are:  zz = 0.06, 0.15, 0.3, 0.5, 0.8, 1 and 1.2. The nondeformed sample,  zz = 0, is taken into account in this study as the material microstructure reference. For each test, the sample is unloaded and the evolution of the axial true strain is followed during 3 hours. Addiego et al. [27], Martin [31] and Ponçot [32] show that after 3 h of recovery at zero stress, the axial true strain keeps nearly stable. The residual true strain,  zz,r , is recorded at the end of this period.

Non-deformed and recovered samples are analysed by WAXS, using the Debye-Scherrer method [2,3]. In the case of uniaxial deformation, many authors have identified a "fibre texture" oriented towards the z-tensile direction (cylindrical symmetry) [35][36][37]. This presupposes that analysis of one 2D-WAXS pattern is sufficient to describe the 3D-spatial arrangement of the crystallites within the polymer. Before WAXS analysis, thickness of samples is reduced to 1 mm by polishing. The analyzed zone is the RVE. Diffraction system and experimental conditions used here are described in details by Addiego et al. [Addiego].

On each pattern, an absorption correction is brought to the intensity magnitude, I, by the Lambert's law, I = I 0 exp(-µt/cos(2)), where 2 is the diffraction angle, µ the absorption coefficient equal in this case to 3.85 cm -1 and t = 1 mm the sample thickness [27,32].

Average orientation of iPP crystallites is quantified by the Hermans'orientation functions, f x/z , which describes the orientation of a given crystallographic axis x relative to one reference direction, usually the z-axis stretching direction [23]. In this paper, we have considered the specific orientation function f c/z since c-axis of the -iPP unit cell is also parallel to the helical axis of the macromolecular chains. Since no (00l) pure c-axis reflections are present in the -iPP WAXS pattern, f c/z is calculated by the combination of two reflecting crystal planes e.g., the (040) and ( 110) planes [Wilchinsly et dahoun].

(

) 1
The terms a = 0.665 nm and b = 2.096 nm are two of the three -iPP unit cell parameters and  = 99 degrees, the angle formed between a and c-axis [40]. The function f c/z varies from a value +1 for c-axis oriented parallel to the z stretching direction, through a value 0 for randomly oriented crystals, to a value -0.5 for c-axis oriented perpendicular to the z-stretching direction.

Raman scattering experiments are performed on non-deformed and post mortem deformed samples using an Aramis  microspectrometer (Horiba Jobin Yvon, Lille, France). A diode laser emitting a 785 nm exciting line is used and allows a substantial reduction in fluorescence. When the laser beam is focused onto the sample, the illumination power density measured on the sample is kept between 15 and 20 mW.µm -2 . All spectra are recorded in backscattered geometry using a 1200 grooves.mm -1 as the diffraction grating and a spectral resolution of about 0.6 cm -1 is achieved. Previously, the Raman shift was calibrated with the strong peak at 521 cm -1 of silicon. Both the incident and scattered lights are collected through an Olympus confocal microscope (100 objective lens, 1000 total magnification, Numerical

Aperture NA = 0.9). Rayleigh scattered light is blocked by using a holographic Notch filter. A half-wave plate is installed on the apparatus allowing the polarization direction of the incident laser radiation to be control. A polarizer is also placed on the backscattering optical path to polarize the Raman scattering radiation. Acquisition time is equal to 20 s and final spectrum is an average of ten spectra. A self-written fitting least-square procedure based on the Levenberg-Marquardt method is used to process spectral data [33]. Residual fluorescence background is eliminated by a polynomial baseline cutting procedure. A mixed Gaussian-Lorentz function is used to decompose each scattering band. Final adjustment gives access to the useful features as position, integrated intensity and full width at half maximum [34].

In situ measurements

The main objective of in situ measurements is to apply the Raman measurement method during the entire course of a tensile test. For this purpose a system coupling a Raman spectrometer and the VidéoTraction TM system is featured (figure 3). Specimens are also machined out from the injection-moulded plates as indicated in figure 1. Contrary to post mortem experiments, thickness of samples is considerably reduced to 500 µm (by polishing).

Main reason is to minimize the deflection of the necking zone ensuring a suitable focalisation of the Raman laser onto the sample during the tensile test. On the same occasion, oriented superficial layers induced by the processing conditions are eliminated.

Raman spectra are collected in backscattered geometry using an iHR250  microspectrometer (Horiba Jobin Yvon, Lille, France) linked to a remote head of probe by means of two optic fibres. The head of probe, placed in face to face to the video-controlled system is equipped with an Olympus long working distance objective lens (12 mm focal).

Laser radiation coming from a diode laser ( = 785 nm) is maintained focused onto the predefined RVE zone where necking appears. A linear polarizer and analyser are also installed in the head of probe. They allow polarizing the incident laser beam and the Raman scattering respectively. Rayleigh scattered light is eliminated through a holographic Notch filter placed on the light path before the diffraction grating, 1200 grooves.mm -1 . The spectral resolution achieved is about 0.6 cm -1 . Conditions of spectral treatment are quite identical as post mortem measurements.

The sample is deformed at 30 °C until its rupture at  zz = 1.6 under a constant true strain rate, zz = 5.10 -3 s -1 . Contrary to post mortem analyses, acquisition time is considerably reduced to 5 s. Mean reason is to minimize structural changes which take place during this period of time. It allows recording 40 spectra during the tensile test. Thus one spectrum corresponds to a strain variation  zz of about 0.04. The experiment is repeated 5 times in the same conditions to ensure reproducibility.

RESULTS AND DISCUSSION

Mechanical behaviour

Figure 4 (a) shows the evolution of the true axial stressstrain curve ( zz - zz ) of the iPP under uniaxial tension at 30 °C remaining a constant axial true strain rate . ε zz = 5.10 -3 s -1 .The stress -strain curve, typical of semi-crystalline polymers, is characterized by three different regimes: i) In the elasto-viscoelastic regime,  zz = 0 -0.09, the initial slope of the curve is related to the elastic Young's modulus, noted E (E = 1.7 GPa). The yield point is defined when stress  y zz achieves 36 MPa and passes through a round-off maximum at strain  y zz = 0.09. ii) In the plastic regime,  zz = 0.09 -0.8, stress is firstly marked by a small stress decrease while the deformation localizes in a sharp neck in the RVE. This is then followed by a symmetric propagation of the neck shoulders towards the specimen heads. iii)

In the plastic stress-hardening regime,  zz = 0.8 -1.6, stress increases exponentially from 32 to 75 MPa.

An example of the axial strain evolution during unloading and recovery is presented in figure 4 (b). In this case, the specimen is deformed until the applied strain reaches 0.8 (step denoted 0 to 1), subsequently unloaded at a strain rate opposite to the initial loading strain rate (step 1 to 2) and finally left to recover at zero stress for 3 h (step 2 to 3). The axial strain drops from 0.8 to 0.71 during the unloading period and from 0.71 to 0.52 during the recovery. Thus the residual axial strain is  zz,r = 0.52. Table 1 summarized the residual axial strains recorded for each post mortem deformed sample under consideration. Nevertheless this indicates that a subtle residual crystalline orientation induced by the injection-moulding process can't be excluded [28]. Globally the increase of the orientation function f c/z with respect to the residual true strain is consistent with a gradual alignment of the c-crystallographic axis towards the z-stretching direction.

Post mortem results

Figures 6 (a) shows whole spectra recorded on a non-deformed ( zz,r = 0) and a large deformed ( zz,r = 0.79) samples by using an x(zz)x polarization configuration according to Porto's nomenclature [41]. In this configuration, electric fields of both incident and scattered lights are parallel to the z-stretching direction. Raman scattering is thus excited and collected principally from chemical bonds oriented towards the z-direction.

In situ results

Figure 8 (a) shows whole spectra recorded on non-deformed and large deformed ( zz = 1.6) samples during in situ experiments. Contrary to post mortem measurements, in situ spectra are noisy. This is caused by the important reduction in acquisition time to 5 s. Figure 8 (b) reports evolution of bands at 973 and 998 cm -1 (after spectral adjustment) in function of various strain levels. Intensity of the 973 cm -1 band increases with the increasing strain. Just the opposite is noted for the 998 cm -1 one. This is in agreement with previous post mortem observations. At low strain, intensity variations are weak, become important at medium strain and become again weak at large strain. Moreover, we can note that these two bands also exhibit a peak shifting. Band at 973 cm -1 shifts towards low wavenumbers while the one at 998 cm -1 shifts towards the high wavenumbers. Raman shift is generally interpreted as an alteration of the atomic vibration (change in the inter-atomic distance) induced by mechanical stress applied at the chemical bond scale [24,33,34]. A shift towards low wavenumbers means that bonds are under an extension stress while a shift towards high wavenumbers is related to a compressive one. Thus, skeletal backbones C-C (band at 973 cm -1 ) which are more and more oriented towards the stretching direction, experience a tensile stress.

Reciprocally lateral alkyl groups CH 3 (band at 998 cm -1 ) undergo a compressive stress.

Mechanisms of orientation

are not yet active. At the early stages of stretching, it is just chains of the amorphous interlamellar layers which undergo deformation [39]. It involves disentanglement, extension and alignment of tie molecules between two adjacent crystalline lamellae. Deformation of amorphous phase is always at least partially reversible [40]. In equatorial regions of spherulites, lamellae are separated and they are subject to a bending solicitation as Nitta et al. proposed it [41]. Furthermore, the existence of a local compressive strain in polar regions induces a lamellar merger and also a lamellar bending. The deformation in diagonal regions is caused by amorphous chains extension and inter-lamellar shearing [42].

ii) In the plastic domain,  zz = 0.09 -0.8, f c/z increases progressively from 0.05 to 0.48. This is related to the beginning of the crystalline chains orientation towards the tensile axis.

At the yield point, amorphous phase cannot accommodate any more the deformation and this is crystals which are affected in turn. Blocks of crystals in equatorial position, attached by tie molecules in extension, are breaked up into smaller folded chains blocks [8][9][10]. Result is quite identical in polar regions where crystal blocks are at first pulled out of the lamellae and then rotate towards the tensile axis. In diagonal regions, the main effect of chain slip is to tilt macromolecules inside the crystalline lamellae and to align them more parallel to the stretching direction [43]. We can note that the inflection point of the increasing evolution of the orientation function is localized at  zz = 0.85, f c/z = 0.36. It suggests that the orientation mechanisms of crystalline chains take a predominant place for this particular strain level.

iii) In the last plastic-hardening domain,  zz = 0.8 -1.6, the orientation function continues to increase but its progression becomes milder. Proportion of crystalline chains which experience tilt towards the tensile axis becomes lower, all being already oriented.

Approaching the rupture deformation, orientation function is maximal and quite constant f c/z = 0.74. It suggests that crystalline chains have achieved their maximal orientation level before the material rupture. The initial lamellar structure is thus transformed into a micro-fibrillar structure with both amorphous and crystalline chains oriented in the tensile direction [8][9][10].

Thus, the material exhibits an important macromolecular anisotropy [44].

Figure 11 compare evolutions of the orientation's function, f c/z , determined in post mortem and in situ measurements in function of the true axial strain. We have to precise that the strains considered for post mortem results correspond to the residual true strain recorded after material recovering. For example an iPP deformed at  zz = 1.2 is characterized f c/z = 0.62.

After material recovering at  zz,r = 0.79, f c/z becomes lower, f c/z = 0.35. The noted differences evidence a disorientation of the crystalline chains during the material recovering sequence noted f c/z . This is consistent with the recovery phenomenon which occurs between the deformation stop and the post mortem microstructural analysis [16][17][18]. Indeed evolution of the deformed semi-crystalline microstructure is time-dependent and unstable. After unloading, the oriented crystal blocks tend to recover their initial random orientation due to the elastic response of the extended tie molecules in the amorphous interlamellae layers. The molecular model presented in figure 12 illustrates the strain recovering associated to the crystalline disorientation which occurs during the material recovery. Figure 13 presents the evolution of the crystalline disorientation f c/z with respect to the strain level  zz achieved in load for post mortem samples. Disorientation becomes more and more pronounced with the strain because the proportion of crystals blocks affected by the recovering mechanism is higher. This result evidences the main advantage of performing in situ experiments namely the measured orientation level corresponds exclusively to the uniaxial deformation, "parasite" contribution of the recovery being avoided.

CONCLUSION

Molecular orientation of the crystalline phase of isotactic polypropylene (iPP) was investigated by polarized Raman spectroscopy. In a fixed polarisation configuration, we have established a measurement method based on the intensity sensitivity of specific Raman scattering bands at 973 and 998 cm -1 to molecular orientation. We proposed the ratio of their integrated intensity as an indicator of orientation. Firstly, the reliability of the method was validated by comparing Raman results with those found by Wide Angle X-ray Scattering on post mortem deformed samples (WAXS-Debye-Scherrer method). A good correlation was found. On the basis of this result, method is then applied to follow the crystalline phase orientation in real-time with the deformation. The experimental set-up developed for this purpose combines a VidéoTraction TM -controlled tensile machine and a Raman spectrometer.

It allows matching one true state of the deformation to one microstructural state with less than 5 s time resolution. Our results agree well with general descriptive model of semi-crystalline polymer deformation suggest by the past by Peterlin [8] and Schultz [9]. Results clearly evidence the recovery phenomenon occurring after sample unloading and confirm the use of in-real-time measurements.

Further works should be devoted to a closer examination of the amorphous phase evolution during deformation more precisely the macromolecular disentanglement and elongation during the elastic and visco-elastic regimes. Collection of data both on amorphous and crystalline phase would be give access to a more complete overview of the microstructural mechanisms of deformation and a better understanding of the resistance properties of semi-crystalline polymers.

FIGURES CAPTIONS

Figure 1: Geometry of the tensile test specimen machined out from an injection-molded plate.

Figures 5 (

 5 Figures 5 (a)-(d) show the WAXS patterns obtained on deformed samples ( zz,r = 0,

Figure 6 (

 6 b) shows influence of the deformation on the intensity of two Raman bands at 973 and 998 cm -1 (after spectral adjustement). They are attributed exclusively to macromolecules in the 3 1 helical conformation of the iPP crystalline phase[13][14][15]. The first one is assigned principally to the C-C asymmetric stretching mode of the skeletal backbones while the second is referred to the rocking mode of the CH 3 alkyl groups in lateral position on the chain. Intensity of the 973 cm - 1 band is high for the deformed sample since proportion of C-C bonds oriented towards the ztensile direction is high. Situation is just opposite for the 998 cm -1 since CH 3 alkyl groups are less numerous being oriented in the z-tensile direction. On the basis of this observation, we propose the use of the intensity bands ratio I 973 / I 998 to estimate orientation of the iPP crystalline chains. Method is applied on deformed samples. Figure7 (a) shows evolution of the bands ratio I 973 / I 998 versus the residual true strain  zz . Error bars correspond to the standard deviation calculated from ten spectra recorded on the RVE of five stretched samples. As expected the Raman bands ratio increases with the increasing deformation. It reflects a progressive orientation of the C-C skeletal backbones of the crystalline chains towards the ztensile direction. Its evolution is comparable with the Herman's orientation function, f c/z .

Figure 7 (

 7 Figure 7 (b) attests of a good agreement between the two series of data since a linear

Figure 9

 9 Figure 9 shows both evolution of the orientation function f c/z (derived from Raman

Figure 2 :

 2 Figure 2: Microstructure of injection-molded specimen within the depth. Optical microscopic view under polarized light performed on a microtomed cross-section.

Figure 3 :

 3 Figure 3: General diagram of the VidéoTraction TM system combined with a Raman spectrometer.

Figure 4 :

 4 Figure 4: (a) Evolution of the tensile true stress-strain curve ( zz - zz ) of iPP recorded at 30 °C during monotonous loading at constant true strain 5.10 -3 s -1 . (b) Example of the production of a stretched sample for post mortem analysis, deformed in load at 0.8 (step 0 to 1), unloaded at 0.7 (step 1 to 2) and recovered at 0.52 maintaining zero stress for 3 h (step 2 to 3).

Figure 5 :

 5 Figure 5: Wide Angle X-ray Scattering (WAXS) patterns obtained on iPP samples uniaxially deformed, unloaded and then recovered at different residual true axial strains  zz,r : (a) 0.0, (b) 0.15, (c) 0.52 and (d) 0.79.

Figure 6 :

 6 Figure 6: Influence of deformation on polarized Raman spectra of iPP (x(zz)x), (a) 750 -1450 cm -1 and (b) 950 -1015 cm -1 . Raman analyses are performed on post mortem deformed samples.

Figure 7 :

 7 Figure 7: (a) Superposed evolutions of the orientation function f c/z and the bands ratio I 973 /I 998 in function of residual strain of iPP. (b) Linear regression between f c/z and I 973 /I 998 .

Figure 8 :

 8 Figure 8: Influence of deformation on polarized Raman spectra of iPP (x(zz)x), (a) 750 -1450 cm -1 and (b) 950 -1015 cm -1 . Raman analyses are performed in situ with the deformation.

Figure 9 :

 9 Figure 9: Evolution of the orientation function f c/z in function of strain of iPP. f c/z is calculated from Raman measurements performed in situ with the deformation.

Figure 10 :Figure 11 :

 1011 Figure 10: Schematic illustration of the crystalline phase orientation occurring during the uniaxial deformation of the iPP microstructure. (inspired from Schultz [9] and Friedriech's models [10]) Figure 11: (a) Evolution of the orientation function f c/z in function of strain of iPP. Comparison between post mortem and in situ measurements. (b) Evolution of the crystalline phase disorientation f c/z in iPP relaxed material in function of various strain level achieved in load.

Figure 12 :

 12 Figure 12: Schematic illustration of the crystalline phase disorientation occurring in iPP microstructure during relaxation sequence of a uniaxially stretched sample.

Figure 1 :

 1 Figure 1: Geometry of the tensile test specimen machined out from an injection-molded plate.

Figure 3 :

 3 Figure 3: General diagram of the VidéoTraction TM system combined with a Raman spectrometer.

Figure 4 :

 4 Figure 4: (a) Evolution of the tensile true stress-strain curve ( zz - zz ) of iPP recorded at 30 °C during monotonous loading at constant true strain 5.10 -3 s -1 . (b) Example of the production of a stretched sample for post mortem analysis, deformed in load at 0.8 (step 0 to 1), unloaded at 0.7 (step 1 to 2) and recovered at 0.52 maintaining zero stress for 3 h (step 2 to 3).

Figure 5 :

 5 Figure 5: Wide Angle X-ray Scattering (WAXS) patterns obtained on iPP samples uniaxially deformed, unloaded and then recovered at different residual true axial strains  zz,r : (a) 0.0, (b) 0.15, (c) 0.52 and (d) 0.79.

Figure 7 :

 7 Figure 7: (a) Superposed evolutions of the orientation function f c/z and the bands ratio I 973 /I 998 in function of residual strain of iPP. (b) Linear regression between f c/z and I 973 /I 998 .

Figure 8 :

 8 Figure 8: Influence of deformation on polarized Raman spectra of iPP (x(zz)x), (a) 750 -1450 cm -1 and (b) 950 -1015 cm -1 . Raman analyses are performed in situ with the deformation.

Figure 9 :

 9 Figure 9: Evolution of the orientation function f c/z in function of strain of iPP. f c/z is calculated from Raman measurements performed in situ with the deformation.

Figure 10 :Figure 11 :

 1011 Figure 10: Schematic illustration of the crystalline phase orientation occurring during the uniaxial deformation of the iPP microstructure. (inspired from Schultz [9] and Friedriech's models [10])

Figure 12 :

 12 Figure 12: Schematic illustration of the crystalline phase disorientation occurring in iPP microstructure during relaxation sequence of a uniaxially stretched sample.

Table 1 :

 1 Preparation of post mortem iPP samples uniaxially stretched.  zz is the strain achieved in load just before unloading,  zz,r is the residual strain measured 3 h after unloading.

	 zz	0.06	0.15	0.3	0.5	0.8	1	1.2
	 zz,r	0.008	0.04	0.15	0.31	0.52	0.65	0.79


text:
Bourson

J Martin

S Margueron

M Fontana

M Cochez

P Bourson

Study of the Molecular Orientation Heterogeneity in Polypropylene Injection-Molded Parts by Raman Spectroscopy

Study of the molecular orientation heterogeneity in polypropylene injection-molded parts by Raman spectroscopy

INTRODUCTION

The mechanical response of the semi-crystalline polymers is strongly dependent on the semi-crystalline microstructure which are induced by the rheological processing conditions such as injection melt temperature T inj (° C), cooling mold temperature T w (° C), flow rate Q inj (mm 3 .s -1 ) and mold geometry [1][2][3][4]. Injection-molding is a wellknown processing method to produce plastic parts for engineering applications. During the process, a hot polymer melt is forced to flow into the cavity of a mold under complex thermo-mechanical conditions characterized by a high cooling rate and an important shear field, especially at the walls of the mold. Thus, the final semicrystalline microstructure of the injection-molded plastic parts is far from being homogeneous and isotropic. This microstructure is widely reported in the literature and shows a hierarchy of macromolecular arrangements and a through-the-thickness morphological gradient [5,6]. In a simple description it consists of a multilayered structure, e.g., skin and core, parallel to the surface of the mold. The skin morphology exhibits a highly-oriented shish-kebab like structure along the injection direction while radial spherulites are observed in the core region [7]. Individual properties and relative importance of each layer determine the mechanical end-use properties of the molded parts [8][9][10]. Few numerical studies tempt to simulate the solidification and the resulting mechanical properties of injection-molded parts [11,12]. Nowadays, there is an increasing demand by plastic transformers of new developments in micro-structural control methods in order to monitor processes in realtime with insurance of quality, reproducibility and rapidity [START_REF] Coates | Polymer processing: in-process measurements[END_REF][START_REF] Lamberti | [END_REF]. Typically, analytical techniques such as light diffraction [START_REF] Schultz | Polymer materials science[END_REF], Differential Scanning Calorimetry (DSC) [START_REF] Gedde | Polymer physics[END_REF], X-ray scattering [START_REF] Turner-Jones | [END_REF], Scanning Electron Microscopy [18], Infra-Red absorption [19] and Nuclear Magnetic Resonance [20] are suitable for polymer morphological analysis but require long and destructive sampling procedure and provide only particular morphological informations. In this paper we propose the use of the Raman spectroscopy to characterize the molecular orientation of injection-molded isotactic polypropylene (iPP) parts.

Raman spectroscopy is a well-suitable technique to study the vibrational characteristics of polymer molecules. This technique is an ideal tool for polymer analysis especially for the study of saturated and unsaturated C-C bonds [START_REF] Hollas | Modern spectroscopy[END_REF][START_REF] Koenig | [END_REF] and provides the analysis of molecular species at different scales, from the chemical bond up to the conformational architecture. The Raman spectrum of iPP has been the object of several investigations, and the assignment of the Raman bands are rather well established [23][24][25]. The development of confocal technique associated with classical Raman spectrometer has considerably improved spatial resolutions in particular axial resolution. It consists on placing a variable aperture at the focal image plan in order to analyze only the light scattered from the interaction focal volume [26].

In this paper, using the polarized Raman micro-spectroscopy, we characterize the molecular orientation heterogeneity in injection-molded iPP parts. We particularly point out the skin-core laminated structure and show how it may be affected with rheological conditions induced by mold geometry variations.

EXPERIMENTAL PROCEDURES

Injection Molded Parts

The iPP resin granules manufactured by Repsol YPF were used to prepare samples. Four rectangular plates were molded by injection using an Arburg 250-75 machine under T inj = 200 °C as the heating temperature, T w = 25 °C as the cooling temperature and P = 10 bars as the pressure. As shown in the Fig. 1, the specimen length L and width w were kept constant to 32 and 16 mm respectively. The thicknesses of the plates were h 1 = 0.5, h 2 = 1, h 3 = 2 and h 4 = 4 mm. The injection direction is parallel to the width w along Z-axis.

Especially for molecular orientation assignment, one specific rectangular dumbbell-shape test specimens with overall dimensions of 34 x 10 x 4 mm 3 was also prepared by injection molding and then uniaxially stretched. The median region is reduced to 9 mm in order to localize necking in this predetermined zone. Nominal strain was calculated from the ratio of the increment of the length between clamps to the initial gauge length. Nominal tensile stress was determined from dividing the tensile load by the initial cross section. Mechanical tests were performed at 25 °C under a constant cross-head speed of about 5 mm.min -1 . A nominal deformation of  N = 200 % is achieved in order to enhance molecular orientation along the stretching direction. The injected material has an average crystallinity index of 53-55 %, as determined by DSC (Perkin-Elmer DSC4 calorimeter) with a major proportion of the monoclinic -crystal form (95 % in weight). The influence of the hexagonal -crystal form is not taking into account in this work. The iPP used has a melting point T m near to 165 °C and a crystallization point T c near to 143 °C.

Raman Spectroscopy

All Raman scattering measurements are carried out on a LabRam microspectrometer developed by Horiba Jobin Yvon (Lille, France). A Helium-Neon laser emitting a 633 nm exciting line is used. The system is equipped with a liquid nitrogen cooled, 2000 x 256 pixels CCD detector (Synapse, Horiba Jobin Yvon). All spectra are recorded in back-scattered geometry using an 1800 grooves.mm -1 centered at 950 cm -1 as the diffraction grating. A spectral resolution of about 1.2 cm -1 is achieved. The useful spectrum under consideration is ranging from 750 cm -1 to 1100 cm -1 . Both the incident and scattered lights are collected through an Olympus confocal microscope (x 100 objective lens, total magnification x 1000, numerical aperture NA = 0.9). In this configuration, the lateral (along x and y-direction) and the axial (along z-direction)

space resolution are about of 1.1 µm and 1,9 µm respectively. When the laser beam is focus onto the sample, the illumination power is about of 16 mW. Rayleigh scattered light is block by using a holographic Notch filter. A set of two polarisers is also installed on the spectrometer and offers the possibility to polarize both the incident and the scattered radiation independently. All Spectra are recorded with the Labspec  acquisition software developed by Horiba Jobin Yvon (Lille, France). Acquisition time is equal to 20 s and final spectrum is an average of ten spectra. A personal fitting leastsquare procedure based on a Lorentzian function adjustment is used to process spectral data and decompose the bands into their wavenumber position, absolute intensity, fullwidth at half maximum and integrated intensity. In this study, we just take into account of the integrated intensity.

Molecular Orientation Assignments

In Fig. 2, we define a coordinate system referred to the sample (O, X, Y, Z). Spectra were recorded on the necked area of the uniaxially drawn dumbbell-shape test specimen. According to Porto's nomenclature [27], two perpendicular polarization geometries were used: X(ZZ)X and X(YY)X respectively parallel and perpendicular to the stretching direction. To explore these two polarization configurations, we used a computer-controlled rotational device permitting to turn the sample by remaining one polarization configuration constant.

Skin Layer Profiles

To characterize molecular orientation heterogeneity through the thickness in the rectangular plate specimens, two different experimental methods were employed. An edge profile was undertaken on the cross-section of one plate specimen from the surface to -600 µm distance. The cross-sections were obtained by carefully cutting plates into two equal parts along specimen width w and then manually polishing with grinding paper of sequentially lower grain sizes. Then an in-depth profile through specimen thickness from the surface to -600 µm depth was performed with a confocal microscope system which is coupled with our Raman spectrometer. The pinhole was adjusted to 100 µm diameter which offers a sufficient axial resolution up to -600 µm in depth. For both experimental methods used, the laser radiation polarization was maintained parallel to injection direction along Z-axis. In order to correlate our measurements, optical microscopy using polarized visible radiation is applied on microtomic slices. Measure of skin thickness is also possible on the optical micrographs recorded.

RESULTS AND DISCUSSION

Molecular Orientation Assignments

The two different scattering configurations defined previously allow the determination of the Raman bands, which are sensitive to chains orientation. Two spectral regions are particularly examined and carefully fitted to a Lorentzian profile.

The first range around 820 cm -1 is structured into three separated bands located at 809, 835, and 841 cm -1 . The 809 and 841 cm -1 correspond to the C-C skeletal symmetric stretching mode v s (C-C) and to the CH 2 lateral alkyl rocking mode r(CH 2 ) respectively [23][24][25]. Both are referred to the 3 1 helical chain conformation or crystalline phase. It is noted that the 835 cm -1 band related to the C-C skeletal symmetric stretching mode v s (C-C) in non-helical chain conformation is not considered in the present discussion.

The second range around 980 cm -1 shows two separate bands at 973 and 998 cm -1 . The 973 cm -1 band has been assigned to the C-C skeletal asymmetric stretching mode v as (C-C) whereas the 998 cm -1 band has been referred to the CH 3 lateral alkyl rocking mode r(CH 3 ). Both are also referred to the 3 1 helical chain conformation or crystalline phase.

The vibrational assignments of the Raman bands considered are given in Table 1.

Fig. 3 (a) and (b) shows the X(YY)X and X(ZZ)X spectra recorded on an unstretched ( N =0 %) and a stretched specimen ( N =200 %) for two wavenumber ranges. By remaining laser light polarization parallel to the tensile direction (X(ZZ)X), bands at 809 cm -1 and 973 cm -1 are dominant while bands located at 841 cm -1 and 998 cm -1 have a weak intensity. This is amplified with the pre-oriented stretched specimen. On the contrary, by remaining laser light polarization perpendicular to the tensile direction (X(YY)X), bands at 809 cm -1 and 973 cm -1 are weak while bands at 841 cm -1 and 998 cm -1 have a dominant intensity. This is also amplified with the preoriented stretched specimen.

These observations can be understood by means of the molecular vibration associated to each Raman band. In the uniaxially pre-oriented stretched sample, skeletal backbone C-C are preferentially oriented parallel to the stretching direction whereas lateral alkyl groups are preferentially oriented perpendicular to the stretching direction. Thus, when the polarization is parallel to the stretching direction, Raman scattering comes principally from C-C backbones which explains the high intensity of bands at 809 and 973 cm -1 . On the contrary, when polarization is kept perpendicular to the stretching direction, Raman scattering comes principally from lateral alkyl groups CH 2 and CH 3 which explains the high intensity of bands at 841 and 998 cm -1 . The use of the two ratio J(809)/J(841) and J(973)/J(998) is therefore a valuable tool to characterize molecular orientation heterogeneity induced in iPP injection-molded parts. Table 2 resumes the ratio for the unstretched and the pre-oriented stretched specimens taken in two perpendicular laser light polarizations.

Edge and In-Depth Profiles

Edge profiles were performed on the cross-sections of the injection-molded rectangular plates. Keeping the laser light polarization toward the injection direction Z, the ratio of integrated intensity defined previously J(809)/J(841) and J(973)/J(998) are used to discriminate preferential molecular orientation in the skin layer. Theses ratio are expected to predominate in the skin region as in the core region. Fig. 4 shows the edge and in-depth Raman profiles. Any ratio variation can be observed for the h 1 = 0.5 mm plate thickness. Due to the fact that the thickness is very thin, the entire plate is highly oriented and can be considered as a skin layer. Fig. 7 shows that the skin layer thicknesses measured by the edge and the in-depth method are in good agreement with optical microscopy experiments. This result proves the usefulness of confocal depth analysis from 0 to -600 µm in estimating skin-core laminated structure induced in injection-molded parts.

Rheological Aspect

From Table 3, it is observed that the skin layer thickness increases with the decreasing of the plate thickness. Fig. 8 illustrates this behavior in representing the skin layer thicknesses e skin previously determined in function of the thickness h of the plates used. It is worth to mention that the plate thickness h correspond also to the thickness of the mold cavity. Thus the skin thickness varies with the mold geometry as 1/h 2 . A comparison can be made with the well-known rheological model which described the shear strain rate  wall at the wall of the mold for a non-Newtonian flow through a rectangular channel [START_REF] Baird | Polymer processing[END_REF][START_REF] Zhu | [END_REF]

: 2 . . 2 h w Q inj wall   (1)
Equation shows that shear strain rate  wall at the wall of the mold is proportional to the flow rate Q inj and inversely proportional to the width w and the thickness h of the mold cavity. Especially,  wall and e skin vary similarly with the thickness of the mold cavity as 1/h 2 as it can be seen in Fig. 8. In other words, skin layer thickness is directly related to the shear strain rate at the wall which depends on the mold geometry. Thin plates experience much stronger melt-shear during the injectionmolding process. It must be pointed out that for a certain sample thickness, the skin layer becomes higher than the sample thickness and we can consider that the entire sample is highly oriented.

CONCLUSIONS

By using two perpendicular laser light polarizations on a highly oriented sample, we have selected specific Raman bands which are sensitive to the molecular orientation. We show that intensity ratio J(809)/J(841) and J(973)/J(998) can be used as the Raman assignments of the molecular orientation of isotactic polypropylene chains.

Taking into account of this result, polarized Raman spectroscopy coupled with the confocal technique was successfully used to study the molecular orientation heterogeneity into injection-molded iPP parts. Especially, the oriented-skin layer thickness was determined and compared with two different experimental ways. Indepth profiles are well-correlated with edge profiles. For the processing window considered, the mold geometry has a great influence on the imposed shear stress level and on the skin development.

FIGURES CAPTIONS

Figure 1 :

 1 Figure 1: The four rectangular injection-molded plates considered in this study and

Figure 2 :

 2 Figure 2: Raman scattering geometry used on a pre-oriented stretched dumbbell-shape

Figure 3 :

 3 Figure 3 : X(ZZ)X and X(YY)X Raman spectra recorded on an unstretched ( N =0%)

Figure 4 :

 4 Figure 4 : Evolution of the ratio J(809)/J(841) (a) and J(973)/J(998) (b) versus depth

Figure 5 :Figure 6 :

 56 Figure 5 : Optical micrograph performed on the cross-section of the thickest plate (h 4 =

Figure 7 :

 7 Figure 7 : Correlation between the skin thicknesses determined by the in-depth and the

Figure 8 :

 8 Figure 8 : Influence of the mold geometry on the skin thickness.

Figure 1 :

 1 Figure 1 : The four rectangular injection-molded plates considered in this study and characterized by different thicknesses h: h 1 = 0.5, h 2 = 1, h 3 = 2 and h 4 = 4 mm.

Figure 2 :

 2 Figure 2: Raman scattering geometry used on a pre-oriented stretched dumbbell-shape

Figure 3 :

 3 Figure 3 : X(ZZ)X and X(YY)X Raman spectra recorded on an unstretched ( N =0%)

Figure 4 :

 4 Figure 4 : Evolution of the ratio J(809)/J(841) (a) and J(973)/J(998) (b) versus depth

Figure 5 :Figure 6 :

 56 Figure 5 : Optical micrograph performed on the cross-section of the thickest plate (h 4 =

Figure 7 :

 7 Figure 7 : Correlation between the skin thicknesses determined by the in-depth and the

Figure 8 :

 8 Figure 8 : Influence of the mold geometry on the skin thickness.

TABLE 1 .

 1 Vibrational assignments of the bands considered between 750 cm -1 and 1100 cm -1 . Ref.[23]

	Position (cm-1)	Vibrational mode	Conformation	Phase
	809	 s (C-C)	Helical	Crystalline
	835	 s (C-C)	Non-Helical	Amorphous
	841	r(CH2)	Helical	Crystalline
	973	 as (C-C)	Helical	Crystalline
	998	r(CH3)	Helical	Crystalline

 s , symmetric stretching;  as , asymmetric stretching; r, rocking

TABLE 2

 2

		. Ratio J(809)/J(841) and J(973)/J(998) calculated for an unstretched ( N =0%)
	and a pre-oriented stretched dumbbell-shape test specimen ( N =200%) in two
	perpendicular light polarizations X(ZZ)X and X(YY)X.
		J(809)/J(841)		J(973)/J(998)
	 N (%)	X(ZZ)X	X(YY)X	X(ZZ)X	X(YY)X
	0	3.8	0.64	7.2	0.61
	200	9.1	0.46	19	0.43

TABLE 3 .

 3 Skin thickness e skin determined by edge and in-depth Raman profiles for the four rectangular plates.

	Sample	Thickness h (mm)	e skin (µm) Edge measurements In-depth measurements Microscopy measurements
	h 1	0.5	500	500	500
	h 2	1	490	480	490
	h 3	2	350	320	340
	h 4	4	160	150	155

ACKNOWLEDGMENTS

The authors are grateful to the engineering researchers from Plate Forme de Technologie Plastinnov (Université Paul-Verlaine de Metz) for valuable help and stimulating discussions all along this study.

evolution of the ratio along the edge profile only for the plate with a thickness of h 4 = 4 mm. As expected, the ratios are much stronger on close surface as in core region and continuously decrease toward the core region. In the core region, the ratios show any variation. Thus the molecular orientation is anisotropic in the skin region. Chains are preferentially oriented along the injection direction and refer to an oriented shish-kebab structure. On the contrary, in the core region the molecular orientation is isotropic and refers to a spherulitic structure [7][8][9][10]. The decrease of the ratio observed between the extreme surface and the core region is related to an increase of the spherulitic proportion from the surface to the core region. A skin thickness e skin can be directly determined from the Fig. 4. Thus a skin layer thickness can be evidenced as e h 4 skin = 160 µm. Fig. 5 presents an optical micrograph obtained from the cross-section of the thickest plate (h 4 =4mm) and skin thickness found has a close similarity as the Raman measurements. Identical investigations were performed on the others rectangular plates and the skin thicknesses determined for each plate are summarized in Table 3.

In-depth profiles were performed through the thickness with maintaining the laser light polarization toward the injection direction along Z-axis. Fig. 6 shows the evolution of the ratio J(973)/J(998) through the thickness for the four different injection-molded plates considered in this study. It is noted that the ratio J(809)/J(841) exhibits the same variations. As remarked previously, close to the surface, the ratio is larger than in internal region. It corresponds respectively to the skin and core regions.

From the Fig. 6, it is also possible to determine a skin layer thickness e skin which is arbitrary limited by the point of inflection of the skin-to-core transition. Table 3 summarizes the skin layer thicknesses e skin versus plate thicknesses h determined from


text:
J Martin

P Bourson

A Dahoun

J M Hiver

The -spherulite morphology of isotactic polypropylene investigated by Raman spectroscopy

Keywords: Raman spectroscopy, Isotactic polypropylene, Crystalline polymorphs, Lamellae orientation

Raman spectroscopy is used to investigate the and -crystalline polymorphs of isotactic polypropylene. Raman spectra of the polymorphs show some minor differences, in particular the wavenumber shifting of certain scattering bands, which reflect the environment of an isolated macromolecular chain within the packing unit cell of the different crystal forms. For example, the shifting of the 842 cm -1 scattering band, which the most susceptible band to cause intermolecular interactions, is used to generate a structural map of one -spherulite among -spherulites with a well-defined resolution. Anisotropy of the crystalline phase orientation within a -spherulite is also investigated by performing polarized Raman spectroscopy measurements. Differences in the polarized spectra are then finely exploited to map the radial lamellae orientation distribution within one -spherulite.

Introduction

Crystalline morphology of semi-crystalline polymers controls the ultimate properties of the material. [START_REF] Ward | Mechanical Properties of Solid Polymer[END_REF] Polymorphism is one of the most important aspects of it. Isotactic polypropylene (noted IPP) have three different crystalline polymorphs referred to as the -, -, and one liquid crystal form named smectic phase. [START_REF] Cheng | Crystalline structures of polypropylene homo-and copolymers[END_REF][START_REF] Lotz | [END_REF][4] All have been identified by X-ray scattering. Crystallization of each depends strongly on the thermo-mechanical conditions of cooling. [START_REF] Varga | Crystalline, melting and supermolecular structure of isotactic polypropylene[END_REF] The predominant -form presents a monoclinic unit cell structure and is principally obtained under typical industrial conditions. [START_REF] Natta | [END_REF]7 The -form shows an hexagonal unit cell structure and its appearance is favored under conditions of shear stress and the use of nucleating agents. [START_REF] Lotz | [END_REF]8,9 The and smetic forms are identified as a triclinic and a pseudo-hexagonal unit cell structure respectively and have a minor occurrence. [10][11][12] The individual chains in all the crystalline forms have a 3 1 helical conformation with three monomer units per helix.

Vibrational analysis of semi-crystalline polymer, such as InfraRed (IR) or Raman spectroscopy, are particularly well-adapted for the characterization of the overall chain structure regularity because IR absorption and Raman scattering spectra are sensitive to intra-molecular interactions. 13 For example, Snyder et al. 14 shows that it is possible to distinguish unambiguously the three isomeric configurations of polypropylene with Raman spectroscopy. Minogianni et al. 15 and Gatos et al. 16 , which exploited the Raman sensibility to IPP 3 1 helical conformation, propose a spectral criterion of the crystallinity degree. On the contrary, it is often assumed that inter-molecular interactions between neighbouring polymer chains in the crystal lattice can be ignored since IR absorption and Raman scattering spectra are insensitive to the lateral order of the crystalline phase. [START_REF] Painter | The Theory of Vibrational Spectroscopy and Its Application to Polymeric Materials[END_REF] Thus, a handful of works have compared the different crystalline polymorphs of IPP.

Results collected by Beckett et al. [START_REF] Beckett | [END_REF] and Goldstein et al. 19 on far-IR experiments show absorption band splitting which is discussed in terms of inter-molecular vibrational coupling within the crystal unit cell. Chalmers et al. 20 tend to review Raman band shifting noted for the -, -, and smetic polymorphs of IPP in a large wavenumber range of about 40 -3100 cm -1 . In recent papers Ellis et al. 21,22 have used the minor differences noted between the mid-IR spectra of the and -IPP forms to generate maps of -spherulites. Images obtained have a poor resolution related to the low spatial resolution of IR micro-spectroscopy (10-15 µm). 23 Same authors have performed high-resolution IR micro-spectroscopy experiments using synchrotron IR sources and show an enhancement of the image resolution. 24 By contrast, it is easier to achieve a beam waist of better than 1 µm in Raman micro-spectroscopy by the use of visible laser sources and microscope objective lens.

In this paper we identify differences between the Raman spectra of the and polymorphs of IPP and used them to construct structural maps of the crystalline morphology with high resolution. In a second part, we focus our attention to crystalline anisotropy in such structures.

Experimental

Material

The used IPP was manufactured and processed by Atochem (ref.: 3050 MN1). The isotactic grade has a broad molecular weight distribution, as assessed by gel permeation chromatography, with M w = 75940 g/mol and M n = 262200 g/mol. Pellets of the polymer were subsequently processed by intrusion in a thick rectangular mould (300 x 200 x 15 mm). This technique consists in slowly extruding the melt at a temperature of 230 °C in the mould maintained at 30 °C feeding under the extruding pressure of 6 MPa during the cooling sequence (240 s). Thermo-mechanical conditions of the intrusion process lead to a predominant -spherulitic morphology but also favor -spherulites crystallization in the inner region of the plate as it is seen schematically in Fig. 1.

To make sure of the presence of the -spherulites, external and internal regions are independently characterized by X-ray scattering experiments using the K  radiation of copper ( = 0.1542 nm).

Patterns are recorded at room temperature with a Philips Geiger counter X-ray diffractometer at 1degree per minute over a 2 Bragg's angle range of 10-40 degrees. X-ray patterns of the Fig. 2 show that external regions are only composed of -spherulites since scattering bands located at 2 = 14.12, 16.9, 18.52, 21.28, 21.

Polarized light microscopy

Polarized light microscopy is used both to observe the polymorphic morphology of the intruded isotactic polypropylene material (shape, size and type of spherulites) and to isolate one particular -spherulite for its later Raman mapping. All micro-photographs are captured by a CCD camera surmounting an Olympus microscope (model AX 70). Last one works in light transmission configuration and is equipped of a set of two crossed polarizers allowing study of birefringent crystals such as spherulites of IPP. [START_REF] Wunderlich | Macromolecular Physics: Crystal structure, Morphology, Defects[END_REF][START_REF] Samuels | [END_REF] Samples observed consist in thin slices of about 30 µm cutted out from the intruded plates as the Fig. 1 shows it. A rotary microtome, developed by Leica (model JUNG RM 2055), is used for sectioning the plate. Slices are then carefully disposed between two cover glasses before their observation.

Raman micro-spectroscopy

All Raman scattering measurements are carried out on an Aramis  micro-spectrometer developed by Horiba Jobin Yvon (Lille, France). A diode laser emitting a 785 nm exciting line is used and allows a substantial reduction in fluorescence. The system is equipped with a liquid nitrogen cooled, 2000 x 256 pixels CCD detector (Synapse, Horiba Jobin Yvon). All spectra are recorded in back-scattered geometry using a 1200 grooves.mm -1 as the diffraction grating and a spectral resolution of 0.6 cm -1 is achieved. Raman spectra were calibrated using the neon emission spectrum. The useful spectrum under consideration is ranging from 750 cm -1 to 1550 cm -1 . Both the incident and scattered lights are collected through an Olympus confocal microscope (x 100 dry objective lens, total magnification x 1000, numerical aperture NA = 0.9 and hole aperture 80 µm).

In this configuration, the lateral and the axial space resolution is about 0.8 µm and 1 µm respectively. When the laser beam is focus onto the sample, the illumination power is ranged between 15 and 20 mW. Rayleigh scattered light is blocked by using a holographic Notch filter.

A set of two polarisers is also installed on the spectrometer and offers the possibility to polarize both the incident and the scattered radiation independently. All Spectra are recorded with the Labspec  acquisition software developed by Horiba Jobin Yvon (Lille, France). A personal fitting least-square procedure based on a Lorentzian function adjustment is used to process spectral data and gives access to the useful position, full width at half maximum and integrated intensity of peaks.

Mapping of the crystalline morphology is performed on a selected area of 185 µm x 180 µm using a point-by-point scanning mode with a step size of 1 µm. Acquisition time is about 1 s which correspond to a total measurement time around 10 hours for 33300 spectra individually collected.

Two mapping are performed:

-A first one is carried out without polarization of the incident and scattered radiation.

-A second one is performed by maintaining constant polarization geometry, marked VV, where the polarization direction of the incident radiation is parallel to the polarization direction of the scattered radiation.

Results and discussion

Differentiation of and -polymorphs

Fig. 3 shows a polarised optical microphotograph of an area containing both and -spherulites of IPP. The bright structure in the centre of the microphotograph corresponds to a -spherulite because hexagonal structure of iPP present a high and a negative birefringence. [START_REF] Wunderlich | Macromolecular Physics: Crystal structure, Morphology, Defects[END_REF][START_REF] Samuels | [END_REF] This -spherulite is surrounded by many darker -spherulites which is related to the biaxial lamellae orientation within the monoclinic structure. 27 Average diameters of the and -spherulites are around 30 µm and 60 µm respectively. Fig. 4 shows Raman spectra consecutively recorded on a and -spherulite in the range of 750 -1500 cm -1 . Spectra are very similar to each other and differences are subtle. Only very small wavenumber shifts in a number of Raman scattering bands can be observed. Intensity changes in the spectra are not taking into account in this paper. Band shifting is particularly illustrated in Fig. 5 for one Raman band located at 842 cm -1 and principally assigned to the rocking mode of the CH 2 alkyl group. 13,14 A shift of about 1 cm -1 is noted between the and the -spherulite. The shift noted for all bands (Tab. 1) does not exceed 1 cm -1 . This is not surprising because the 3 1 helical conformation of the iPP chains in each crystal remains identical leading to identical intra-molecular interaction energies. Spectra only differ in that helices have not the same lateral steric environment which influences the vibrational motions of the chemical bonds. All the bands which exhibit a shift are assigned in part with a specific Raman vibrational mode of the lateral alkyl groups (CH 2 , CH 3 ) of the IPP monomer. 13,14 Authors such as Beckett and al. [START_REF] Beckett | [END_REF] and Goldstein and al. 19 ,who used principally IR spectroscopy, confirm that the pendant methyl group CH 3 is clearly the most susceptible to generate inter-molecular interaction.

One can safely assume that the observed wavenumber shift of the Raman bands are due to the crystallographic changes between a -IPP monoclinic and a -IPP hexagonal unit cell rather than thermo-mechanical effects during the crystallization sequence. Eventual compressive (and / or extensional) residual stresses at the spherulite boundaries which could appear during the sequential growth are relaxed by the amorphous phase. [START_REF] Treloar | The Physics of Rubber Elasticity[END_REF] This is comforted by Gatos et al. 16 who show a higher amount of amorphous phase at the circumference of the -spherulites than in their centre.

On the basis of the results obtained, attention is focused on the shift of the 842 cm -1 band. This is followed over an area of 185 µm x 180 µm delimited by the black box, marked A, on the microphotograph of the Fig. 3. Experimental conditions of the mapping are previously described in the experimental part of the study. Fig. 6 presents the two-dimensional evolution of the 842 cm -1 band shifting over the considered area. As expected this evolution generates an image representative of the crystalline morphology of two polymorphs of IPP. It is noted that the correlation between the polarised optical micrograph and the Raman structural map is good because shape and size of the -spherulite are respected. Contour of the -spherulite is well-marked showing its characteristic concave edges. Similar correlations can be made with other bands which exhibit a band shifting as noted in Tab. 1.

Morphological anisotropy within a -spherulite

As the Fig. 7 shows it, many spectra are recorded at fixed distance and at variable angle  from the nucleus of one -spherulite. Measurements are performed using polarization of both the incident and the scattered radiation. The polarization direction of the incident radiation is parallel to the polarization direction of the scattered one (configuration marked VV). Fig. 8 shows the Raman spectra recorded at two extreme angular positions  = 0 and 90 degrees in the range of 960 -1010 cm -1 . Intensity of the bands located at 973 cm -1 and 998 cm -1 is modified as a function of the angular position. The first band is assigned principally to the C-C asymmetric stretching mode of the skeletal backbones while the second is referred to the rocking mode of the CH 3 lateral alkyl groups. 13,14 For  = 0 degree, it is seen that band at 973 cm -1 is dominant while band at 998 cm -1 is weak. For  = 90 degrees, the band at 973 cm -1 has a minor intensity with respect to the band at 998 cm -1 . Observations have to be related to the crystalline lamellae orientation within a -spherulite. [START_REF] Wunderlich | Macromolecular Physics: Crystal structure, Morphology, Defects[END_REF][START_REF] Samuels | [END_REF][START_REF] Schultz | Polymer Materials Science[END_REF][START_REF] Aboulfaraj | [END_REF] As the schematic diagram of the Fig. 9 shows it, c-axis of the lamellae at  = 0 degree (equatorial region) is parallel to the direction of polarization VV. In this case, Raman scattering is excited principally from the C-C skeletal backbones explaining the high value of the 973 cm -1 band intensity. [START_REF] Bower | [END_REF][30][31][32] Simultaneously, Raman scattering produced from the CH 3 lateral alkyl groups is less activated which explains the low value of the 998 cm -1 band intensity. This tendency is completely inversed when the c-axis of the lamellae at  = 90 degrees (polar region) is perpendicular to the direction of polarization. It appears that the ratio I(973) / I(998) of these two bands intensity can be used as a Raman criterion of the lamellae orientation of IPP. Fig. 10 presents this ratio evolution I(973) / I(998) in function of the angular position in the -spherulite considered (Fig 7 ) and a periodic variation is observed confirming the radial orientation distribution of lamellae within a -IPP spherulite. Maxima correspond to lamellae in equatorial region of the -spherulite while minima are attributed to lamellae in polar region of it (Fig. 9). Fig. 11 presents the mapping of a -spherulite by following the evolution of the ratio I(973) / I(998) (black box marked B in Fig. 7). Equatorial, diagonal and polar region are well-delimited. This structural image remind the well-known Maltese-cross extinction obtained on a -spherulite using polarized optical microscopy. [START_REF] Wunderlich | Macromolecular Physics: Crystal structure, Morphology, Defects[END_REF][START_REF] Samuels | [END_REF][START_REF] Schultz | Polymer Materials Science[END_REF][START_REF] Aboulfaraj | [END_REF][35]

Conclusions

The measurements performed in this study demonstrate that Raman spectroscopy allows differentiating the and -crystalline polymorphs of IPP. Explanation of the subtle differences noted between the Raman spectra of and -forms has to be related to the inter-molecular interactions which are different in each crystal lattice. Spectral differences can be exploited to generate a well-resolved image of the crystalline polymorphs. Moreover, Raman spectroscopy in one polarized configuration is applied to develop a spectral criterion of the lamellae orientation in IPP. This criterion is then used to image the crystalline orientation anisotropy within a -spherulite and confirms experimentally the radial distribution of lamellae.

Contrary to other techniques as infra-red spectroscopy and X-ray scattering, acquisition times are reduced and spatial resolution (1 µm) is considerably enhanced.

List of tables

Table 1: Observed Raman frequencies (cm -1 ) for  and  polymorphic forms of IPP

List of figures

  80 and 21.86 degrees are attributed to the -(110), -(040), -(130), -(111), -(131) and -(041) crystallographic planes of the -monoclinic crystal lattice respectively. 8 Apparition of the scattering band located at 2 = 16.2 degrees corresponding to the -(300) crystallographic plane of the -hexagonal crystal lattice indicates the presence of both and -spherulites in the inner region.

Figure 1 :

 1 Figure 1: Schematic diagram of the microtomed slice used in the study and cutted from an IPP

Figure 2 :

 2 Figure 2: X-ray scattering diagrams of the internal (----) and external () region of the IPP

Figure 3 :

 3 Figure 3: Polarized optical microphotograph of  and -IPP spherulites. The box marked A

Figure 4 :

 4 Figure 4: Individual Raman spectra recording on a  and -IPP spherulites, in the range of 750 -1550 cm -1 .

Figure 5 :

 5 Figure 5: Individual Raman spectra recording on a  and -IPP spherulites, in the range of 825 -860 cm -1 .

Figure 6 :

 6 Figure 6: Raman imaging of one -IPP spherulite surrounded of -IPP spherulites. The map (185 µm x 180 µm) is constructed using the 842 cm -1 band shifting.

Figure 7 :

 7 Figure 7: Polarized optical microphotograph of  and -IPP spherulites. Experimental arrangement

Figure 8 :

 8 Figure 8: Polarized Raman spectra recorded at angles  = 0 and 90 degrees on a -IPP spherulite in the range of 960 -1010 cm -1 .

Figure 9 :

 9 Figure 9: Schematic diagram of lamellae orientation distribution within a -IPP spherulite with

Figure 10 :

 10 Figure 10: Ratio I(973) / I(998) of integrated intensity of the 973 cm -1 and 998 cm -1 bands versus

Figure 11 :

 11 Figure 11: Raman imaging of the lamellae orientation anisotropy within one -IPP spherulite. The

Figure 1 :

 1 Figure 1: Schematic diagram of the microtomed slice used in the study and cutted from an IPP

Figure 2 :

 2 Figure 2: X-ray scattering diagrams of the internal (----) and external () region of the IPP

Figure 3 :

 3 Figure 3: Polarized optical microphotograph of  and -IPP spherulites. The box marked A

Figure 4 :Figure 6 :

 46 Figure 4: Individual Raman spectra recording on a  and -IPP spherulites, in the range of 750 -

Figure 7 :

 7 Figure 7: Polarized optical microphotograph of  and -IPP spherulites. Experimental arrangement

Figure 8 :

 8 Figure 8: Polarized Raman spectra recorded at angles  = 0 and 90 degrees on a -IPP spherulite in

Figure 9 :

 9 Figure 9: Schematic diagram of lamellae orientation distribution within a -IPP spherulite with

Figure 10 :

 10 Figure 10: Ratio I(973) / I(998) of integrated intensity of the 973 cm -1 and 998 cm -1 bands versus

Figure 11 :

 11 Figure 11: Raman imaging of the lamellae orientation anisotropy within one -IPP spherulite. The

Table 1 :

 1 Observed Raman frequencies (cm -1 ) for  and  polymorphic forms of IPP

	Polymorph				frequency (cm -1 )
	-IPP	809	835	842	900	940	974	998	1037
	-IPP	810	834	843	901	940,5	974,5	998,5	1037