Stress-Induced Polymorphic Transformations and Mechanical

CRYSTAL GROWTH & DESIGN

Stress-Induced Polymorphic Transformations and Mechanical Properties of Isotactic Propylene-Hexene Copolymers Claudio De Rosa,*,† Finizia Auriemma,† Odda Ruiz de Ballesteros,† Stefania Dello Iacono,† Dario De Luca,† and Luigi Resconi‡,§

2009 VOL. 9, NO. 1 165–176

Dipartimento di Chimica “Paolo Corradini”, UniVersita` di Napoli “Federico II”, Complesso Monte S. Angelo, Via Cintia, I-80126 Napoli, Italy, and Basell Polyolefins, Centro Ricerche G. Natta, P.le G. Donegani 12, I-44100 Ferrara, Italy ReceiVed January 28, 2008; ReVised Manuscript ReceiVed October 22, 2008

ABSTRACT: A study of the relationships between the stress-induced phase transitions and the mechanical properties of isotactic propylene-hexene copolymers with hexene concentration in the range 1-26 mol%, prepared with metallocene catalysts, is reported. Hexene units are included in crystals of R form of isotactic polypropylene (iPP) and produce large disturbance of the crystalline lattice and a consequent decrease of melting temperature, degree of crystallinity, crystallite size, and plastic resistance of the crystals. Defective crystals of R form rapidly transform by stretching into the mesomorphic form of iPP that, in turn, facilitates further stretching up to very high values of deformation of nearly 900-1000%, resulting in high flexibility. This explains the experimental observation that the presence of hexene comonomeric units induces a strong enhancement of ductility, flexibility, and toughness, compared to the highly stereoregular homopolymer prepared with the same catalyst. These copolymers show mechanical properties of highly flexible materials with values of the tensile strength, elastic modulus, and resistance to the plastic deformation that depend on the degree of crystallinity and the occurrence of phase transitions during deformation and can be easily tailored by changing the hexene concentration. Introduction In recent articles, the structure of isotactic propylene-hexene random copolymers (iPPHe), prepared with metallocene catalysts, with hexene concentration variable in a wide range between 2 and 26 mol %, has been described.1-4 It has been observed that the presence of hexene at concentrations higher than 10 mol % produces crystallization of iPPHe copolymers into the trigonal form of isotactic polypropylene (iPP).1-4 Copolymer samples with hexene content lower than 10 mol% crystallize, instead, in the usual R form of iPP.1-4 It has been demonstrated that hexene units are partially included in both crystals of R form and trigonal form,1,2 but the crystallization of the trigonal form at high hexene concentrations allows incorporation of high amounts of hexene units, higher than that in crystals of R form.1-4 The inclusion of hexene units in the crystals induces a suitable increase of density that allows crystallization of threefold helical chains in the trigonal form (space group R3jc),1,2 giving a structure similar to that of form I of isotactic polybutene (iPB).5 This form has never been observed for iPP homopolymer because, in the absence of bulky side groups, it would have a too-low density.1,2 The phase transformation of the R form into the trigonal form with increasing hexene concentration represents a clear example of the influence of the presence of defects on the polymorphic behavior of iPP. Incorporating defects in iPP chains has been extensively used for modifying the iPP properties. In particular, the random copolymerization of propylene with comonomers of different size has been described as an efficient tool to improve the limited impact resistance of iPP.6 For instance, random copolymers of propylene with ethylene and butene have been commercially produced and used mainly as films in the * To whom correspondence should be addressed. Tel.: ++39 081 674346. Fax: ++39 081 674090. E-mail: [email protected]. † Universita` di Napoli “Federico II”. ‡ Basell Polyolefins. § Present address: Borealis Polyolefine GmbH, St. Peter Str. 25, 4021 Linz, Austria.

packaging market, thanks to their good transparency, higher impact strength, and lower heat-seal temperature compared with those of the iPP homopolymer.6 The impact of constitutional defects on material properties depends on type and concentration of the comonomer. It is expected that incorporation of hexene comonomeric units in iPP chains may produce physical and mechanical properties different from those of copolymers with ethylene and butene. Moreover, the physical and mechanical properties depend on the polymorphic transformations that may occur during deformation of the material induced by application of tensile stress. The phase transitions and the crystalline form that develops during deformation, in turn, are strongly affected by the inclusion of comonomeric units in the crystals of iPP, with consequent remarkable influence on the mechanical behavior. In this article, we report a study of the crystallization behavior of iPPHe copolymers in oriented fibers and of the relationships between the structural transformations that occur during stretching and the mechanical properties. The effect of the presence of hexene units on the polymorphic transitions occurring during plastic deformation is investigated. The iPPHe copolymer samples have been prepared with metallocene catalysts, which yield highly stereoregular copolymers, with a very small concentration of stereodefects and regiodefects, and a truly random comonomer distribution. These properties afford an opportunity for studying the effect of hexene units on the physical properties of iPP. Experimental Section Samples of iPPHe copolymers were prepared in liquid monomers or hexane solutions at temperatures between 50 and 70 °C with the C2-symmetric metallocene catalysts shown in Chart 1,7-9 activated with methylalumoxane (MAO). All samples are listed in Table 1. The N in iPPHeN indicates the hexene concentration in mol % in the copolymer. The microstructural data of all samples, hexene concentration, and contents of rr stereodefects and regiodefects were obtained from 13C NMR analysis. All spectra were obtained using a Bruker DPX-400

10.1021/cg800102f CCC: $40.75  2009 American Chemical Society Published on Web 12/04/2008

166 Crystal Growth & Design, Vol. 9, No. 1, 2009

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Chart 1. Structures of Metallocenes Used in This Study

Table 1. Compositions (mol % Hexene), Weight Average Molecular j w), Polydispersities (M j w/M j n), Melting Temperatures of Masses (M Melt-Crystallized Samples (Tm), Contents of Stereoerrors (rr), and Concentrations of Secondary 2,1-Erythro Units (2,1e) of Homopolymer Sample iPPA and Propylene-Hexene Copolymers Prepared with the MAO-Activated Metallocenes of Chart 1 sample a

iPPA iPPHe1.2a iPPHe2.0a iPPHe2.5a iPPHe3.2a iPPHe3.7a iPPHe4.2a iPPHe6.8a iPPHe9.0a iPPHe11.2a iPPHe18b iPPHe26c

mol % hexene

jw M

0 1.2 2.0 2.5 3.2 3.7 4.2 6.8 9.0 11.2 18.0 26.0

237500 699600 122000 430800 152000 333200 291700 239500 209800 266300 217400 184500

j w/M j n Tm (°C)d [rr] (%)e 2,1e (%)f M 2.2 2.3 2.4 2.0 2.0 2.0 2.0 2.2 2.0 1.9 1.9 2.0

152 140 127 125 119 115 109 96 87 70 49 50

0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.2 3.5 0.2

0.8 0.1 0.2 0.2 0.2 0.2 0.2 0.1 0.1 0.1 0 0.1

a Samples prepared with catalyst A of Chart 1. b Sample prepared with a C1-symmetric catalyst10 (metallocene 19 of ref 10a). c Sample prepared with catalyst B of Chart 1. d Melting temperatures of samples crystallized from the melt by cooling the melt to room temperature at 10 °C/min, evaluated from DSC scans at heating rate of 10 °C/min. e [rr] is the percentage of primary stereoerrors over all monomer units, [rr] ) [mrrm] + [mrrr]. It is not determinable and is assumed to be the same as that found in the corresponding homopolymer. f Secondary insertions 2,1 are only of the erythro type, and their amount is normalized over all monomer units.

spectrometer operating in the Fourier transform mode at 120 °C at 100.61 MHz. The samples were dissolved with a 8 wt %/v concentration in 1,1,2,2-tetrachloroethane-d2 at 120 °C. The carbon spectra were acquired with a 90° pulse and 15 s of delay between pulses and CPD (WALTZ 16) to remove 1H-13C coupling. About 1500-3000 transients were stored in 32 K data points using a spectral window of 6000 Hz. The peak of the propylene methine carbon atoms was used as internal reference at 28.83 ppm. The resonances were assigned according to Forlini et al.;11 the 1-hexene content was determined from the diad distribution using the following equations (P ) propylene, H ) hexene):

HH ) 100SRR(HH)/ΣSRR PH ) 100SRR(PH)/ΣSRR PP ) 100SRR(PP)/ΣSRR [P] ) PP + 0.5PH [H] ) HH + 0.5PH The product of reactivity ratios r1 × r2 was determined from the diads according to Kakugo et al.,12 r1 × r2 ) 4[PP] × [HH]/[PH]2. The mass average molecular masses were evaluated from size exclusion cromatography (SEC). Films used for the X-ray diffraction characterization and the analysis of mechanical properties were prepared by compression molding. Powder samples were heated at temperatures higher than the melting temperatures under a press at low pressure and cooled to room temperature. The melting temperatures were obtained with a differential scanning calorimeter (DSC) Perkin-Elmer DSC-7 performing scans in a flowing

N2 atmosphere and heating rate of 10 °C/min. The melting temperatures reported in Table 1 correspond to those of melt-crystallized compression molded samples, and the relative DSC thermograms are shown Figure 4A. They are slightly different from those of the as-polymerized samples reported for the same samples in Table 1 of ref 2. Oriented fibers of the copolymer samples were obtained by stretching at room temperature and at a drawing rate of 10 mm/min compression molded samples at different degrees of deformation, up to values of 500-600%, that is up to final lengths of the specimens 6L0 ÷ 7L0, with initial length L0 of 1 mm. X-ray diffraction patterns (WAXS) were obtained with Ni-filtered Cu KR radiation. The powder profiles were obtained with a Philips diffractometer with continuous scans of the 2θ angle and scanning rate of 0.02 deg/s. The fiber diffraction patterns were recorded on a BASMS imaging plate (Fujifilm) using a cylindrical camera and processed with a digital imaging reader (Fuji BAS 1800). The indices of crystallinity (xc) were evaluated from the X-ray powder diffraction profiles by the ratio between the crystalline diffraction area (Ac) and the total area of the diffraction profile (At), xc ) Ac/At. The crystalline diffraction area was obtained from the total area of the diffraction profile by subtracting the amorphous halo. The procedure used for the evaluation of the amorphous halo for each sample and for the subtraction is the same as that used in ref 2. The scattering of amorphous phases of samples iPPHe18 and iPPHe26 was obtained from the X-ray diffraction profiles of the amorphous samples prepared by cooling the melt to room temperature, before crystallization occurs upon aging at room temperature. The amorphous halo of copolymer samples with hexene contents lower than 9 mol %, which crystallize in the R form of iPP, was obtained from the X-ray diffraction profile of an atactic polypropylene. The scattering of the amorphous phases of samples iPPHe9 and iPPHe11, which crystallize as mixtures of crystals of the R form and the trigonal form, was obtained by the average of the amorphous haloes of atactic polypropylene and of the sample iPPHe26, weighted with respect to the relative amount of the two crystalline phases. The amorphous scattering was then scaled and subtracted to the X-ray diffraction profiles of the semicrystalline samples. The crystallite size was determined from the broadness of reflection peaks in the X-ray powder diffraction profiles of melt-crystallized samples. The half-height widths of the 110 and 040 reflections of the R form at 2θ ≈ 14 and 17°, for samples that crystallize in the R form, and of the (110)T reflection of the trigonal form at 2θ ≈ 10° for the samples iPPHe18 and iPPHe26 were evaluated, and the crystallite size along directions normal to the 110 (L110) and 040 (L040) planes was calculated by the Scherrer equation. Small-angle X-ray scattering (SAXS) data of compression-molded films were collected using a Kratky compact camera SAXSess (Anton Paar) in the slit collimation configuration, attached to a conventional X-ray source (Cu KR, wavelength λ ) 1.5418 Å). The scattered radiation was recorded on a BAS-MS imaging plate (Fujifilm) and processed with a digital imaging reader (Fuji BAS 1800). The range of scattering vectors 0.1 nm-1 e q e 2 nm-1, where q ) (4π sin θ/λ) and 2θ is the scattering angle, was analyzed. After subtraction for dark current, the empty sample holder, and a constant background due to thermal density fluctuations, the slit smeared data were deconvoluted with the primary beam intensity distribution using the SAXSquant 2.0 software to obtain the corresponding pinhole scattering (desmeared) intensity distribution. The constant value of intensity approximating the background Iback was found by fitting the smeared SAXS intensity curve in the range 2 < q < 4 nm-1 with the function:13

Iback + bq-3

(1)

where Iback and b are fitting parameters. Mechanical tests were performed at room temperature on compression-molded films with a miniature mechanical tester apparatus (Minimat, by Rheometrics Scientific), following the standard test method for tensile properties of thin plastic sheeting ASTM D882-83. Compression-molded films were prepared by heating powder samples at temperatures higher than the melting temperatures between perfectly flat brass plates under a press at very low pressure, and slowly cooling to room temperature. Special care was taken to obtain films with uniform thickness (0.3 mm) and minimize surface roughness, according to the recommendation of the standard ASTM D-2292-85. Rectangular specimens 10-mm long, 5-mm wide and 0.3-mm thick were stretched

Phase Transitions in Propylene-Hexene Copolymers

Crystal Growth & Design, Vol. 9, No. 1, 2009 167

Figure 1. 13C NMR spectrum of the copolymer sample iPPHe9.0, with 9.0 mol % of hexene. The resonances assigned to carbon atoms of propylene (P) and hexene (H) units and corresponding to PP, PH, and HH sequences are indicated. The carbon atoms of the butyl branches (4B) belonging to hexene units are indicated with the symbol 4Bn. up to the break or up to a given deformation ε ) [(Lf - L0)/L0] × 100, where L0 and Lf are the initial and final lengths of the specimen, respectively. Two benchmarks were placed on the test specimens and used to measure elongation. In the mechanical tests, the ratio between the drawing rate and the initial length was fixed equal to 0.1 mm/(mm × min) for the measurement of Young’s modulus and 10 mm/(mm × min) for the measurement of stress-strain curves and the determination of the other mechanical properties (stress and strain at break and tension set). The reported values of the mechanical properties were averaged over at least five independent experiments.

Results and Discussion Chain Microstructure (13CNMR). The highly isoselective C2-symmetric metallocenes A7,8 and B9 produce high molecular mass and highly stereoregular iPP homopolymer and iPPHe copolymers containing very small concentrations of stereoerrors (about 0.2 mol % of rr triad defects) and regiodefects due to secondary 2,1-erythro (2,1e) insertions of propylene units (nearly 0.8 mol % in the homopolymer sample iPPA and 0.1-0.2 mol % in the copolymers; Table 1).7-9 Regioinverted hexene units could not be identified by 13C NMR. The 13CNMR spectrum of the sample iPPHe9.0, prepared with catalyst A, showing the assignment of resonances to the hexene branches, is reported in Figure 1, as an example. The SEC curves of all samples show narrow molecular weight distributions, with Mw/Mn ≈ 2, typical of single-center metallocene catalysts (Table 1). Narrow molecular weight distributions, values of product of reactivity ratios r1 × r2 close to 1, and the NMR data indicate that all copolymer samples have a random distribution of comonomers and homogeneous intermolecular composition. Structural (WAXS) Analysis. The X-ray powder diffraction profiles of melt-crystallized compression-molded films of iPPHe copolymers are reported in Figure 2. Samples with hexene content up to 11 mol % crystallize from the melt basically in the R form, as indicated by the presence of the (130)R reflection at 2θ ) 18.6° of the R form and the absence or the very low

Figure 2. X-ray powder diffraction profiles of samples of the homopolymer iPPA (a) and iPPHe copolymers (b-m) with the indicated concentration of hexene units (He), crystallized from the melt by compression molding and cooling the melt to room temperature (or below room temperature) at nearly 10 °C/min. For the samples iPPHe18 and iPPHe26, which are amorphous after cooling and crystallize by aging at room temperature, the diffraction profile of samples aged for two months are reported (l, m). For each sample, the diffraction profile of the amorphous phase is indicated as a dashed line.

intensity of the (117)γ reflection of the γ form at 2θ ) 20.1° in the diffraction profiles of Figure 2b-i. However, the low intensity of the (130)R reflection and the presence in some samples of the (117)γ reflection with very low intensity (Figure 2e) indicate that R/γ disordered modifications intermediate between R and γ forms,14-17 and/or very small amounts of crystals of γ form, were obtained. Samples iPPHe18 and iPPHe26 with 18 and 26 mol % of hexene units, instead, do not crystallize from the melt, and amorphous samples are


De Rosa et al.

Table 2. Elastic Modulus (E), Stress (σy) and Strain (εy) at the Yield Point, Stress (σb) and Strain (εb) at Break, Index of Crystallinity (xc), Crystallite Size along Directions Normal to the 110 (L110) and 040 (L040) Planes, Long Period (L), and Thickness of the Crystalline Lamellae (lc) of Compression-Molded Films of iPPHe Copolymers Prepared with Catalysts of Chart 1 sample

mol % hexene

E (MPa)

εy (%)

σy (MPa)

εb (%)

σb (MPa)

xc (%)

L110 (Å)

L040 (Å)

L (Å)

lc (Å)a

iPPA iPPHe1.2 iPPHe2.0 iPPHe2.5 iPPHe3.2 iPPHe3.7 iPPHe4.2 iPPHe6.8 iPPHe9.0 iPPHe11.2 iPPHe18 iPPHe26

0 1.2 2.0 2.5 3.2 3.7 4.2 6.8 9.0 11.2 18.0 26.0

600 ( 10 363 ( 18 354 ( 20 295 ( 8 307 ( 25 273 ( 11 190 ( 8 126 ( 12 74 ( 5 37 ( 6 13 ( 2 3(1

20 ( 2 18 ( 2 18 ( 4 17 ( 3 23 ( 6 21 ( 2 34 ( 2 41 ( 4 48 ( 4 86 ( 19 90 ( 5

28 ( 2 28 ( 3 23 ( 2 22 ( 2 19 ( 1 18 ( 1 11 ( 1 7(1 4(1 2.5 ( 0.5 1.0 ( 0.5

6(1 906 ( 62 970 ( 50 979 ( 47 850 ( 70 900 ( 65 973 ( 66 873 ( 40 937 ( 40 945 ( 55 860 ( 75 915 ( 60

7(3 24 ( 2 24 ( 2 22 ( 3 21 ( 1 24 ( 1 25 ( 3 23 ( 2 22 ( 3 20 ( 4 7(2 7(2

69 63 57 57 55 53 51 48 39 31 31 24

134 114 114 114 114 114 107 89 73 57 133 133

200 179 179 167 160 160 160 160 134 115

128

88

118

67

108

57

105

50

135 180

42 56

a Roughly determined from the product of the lamellar long period L and the index of crystallinity, obtained from X-ray powder diffraction profiles of Figure 2, as lc ) Lxc/100.

obtained by cooling the melt below room temperature. Amorphous specimens of the samples iPPHe18 and iPPHe26 slowly crystallize in the trigonal form upon aging at room temperature for at least 24 h (Figure 2l,m).2 The absence of γ form in melt-crystallized samples of these copolymers is contrary to most literature data published thus far14-32 that indicate that the crystallization of γ form of iPP is favored by the presence of microstructural defects, that is, stereo- and regio-irregularities14-22 and comonomeric units.23-32 The preferential crystallization of the R form in iPPHe copolymers is related to the preferential inclusion of hexene units in crystals of R form,33 in agreement with results reported recently also on propylene-butene copolymers34 that have demonstrated the importance of the effect of inclusion of defects in crystals of R and γ form in driving crystallization of the two polymorphs of iPP.33,34 The index of crystallinity and the crystallite size along directions normal to the 110 (L110) and 040 (L040) planes of meltcrystallized samples are reported in Table 2. The melting temperature, the index of crystallinity, and the crystallite size L110 are also reported in Figure 3 as a function of hexene concentration. The melting temperature (Figure 3A) and the index of crystallinity (Figure 3B) decrease with increasing hexene concentration, whereas the crystallite size (Figure 3B) decreases compared to the homopolymer sample iPPA already for low hexene content, is nearly constant up to hexene concentration of 4-5 mol %, then decreases for hexene contents higher than 5-6 mol %, and then increases again for hexene concentration higher than 10-15 mol %, when the samples crystallize in the trigonal form (Figure 2). These data are in agreement with the finding that at low concentrations hexene units are partially included in the crystals of R form and act as a lattice defect, producing large disturbance of the crystalline lattice and a consequent decrease of the degree of crystallinity2 and crystallite size along directions normal to the chain axes. For concentrations higher than 11 mol %, larger amounts of hexene units are more easily accommodated in the crystalline lattice of the trigonal form, producing a lower decrease of the degree of crystallinity and melting temperature.2 Accordingly, an increase in the coherent length of the crystallites in directions normal to chain axes (i.e., the lattice directions of the trigonal form) is also observed (Figure 3B). SAXS Analysis. The lamellar thickness of iPPHe random copolymers was determined from SAXS patterns of meltcrystallized compression-molded samples. The Lorentz-corrected SAXS profiles of selected copolymer samples and of the homopolymer and the corresponding DSC melting curves are

Figure 3. Melting temperature Tm (0) (A), index of crystallinity xc (O), and crystallite size along the direction normal to the 110 plane L110 (b) (B) of melt-crystallized samples of iPPHe copolymers of Figure 2 as a function of hexene concentration. For the samples iPPHe18 and iPPHe26 with 18 and 26 mol % of hexene, which are amorphous soon after the compression molding, the data correspond to those of compression-molded samples after crystallization at room temperature for two months in the trigonal form.

shown in Figure 4A,B, respectively. The homopolymer sample iPPA displays first- and second-order peaks at q ≈ 0.5 and 1 nm-1 (curve a of Figure 4A), whereas the iPPHe copolymer samples show only one peak. For the samples with hexene content less than 18 mol %, which basically crystallize in the R form, the peak position first increases with increasing hexene content, from q ≈ 0.56 nm-1 for the sample iPPHe2.0 (curve b of Figure 4A) up to q ≈ 0.63 nm-1 for the sample iPPHe6.8 (curve d of Figure 4A), and then decreases to the value of ∼0.48 nm-1 for the less crystalline sample iPPHe11.2 (curve e of Figure 4A). The intensity of these peaks decreases with increasing hexene concentration, whereas a large increase of the SAXS intensity is observed in the low q regions, especially for the sample iPPHe11.2 (curve e of Figure 4A). Finally, for the sample iPPHe18 that crystallizes in the trigonal form, a neat maximum is apparent at q ≈ 0.34 nm-1 (curve f of Figure 4A). According to AFM analysis performed on some metallocenemade iPPHe copolymer samples having similar compositions,29b the SAXS profiles of Figure 4A may be interpreted in terms of


Figure 4. (A) Lorentz-corrected SAXS intensity profiles measured at room temperature and (B) DSC curves of iPPHe copolymers with the indicated concentration of hexene units (He) crystallized from the melt by compression-molding and cooling the melt to room temperature (or below room temperature) at nearly 10 °C/min. For the sample iPPHe18 with 18 mol % of hexene, which is amorphous soon after the compression molding, the profile corresponds to that of compressionmolded sample after crystallization at room temperature for two months in the trigonal form. In (B), arrows indicate the end-melting temperatures Tmax(lmax) at which the crystals of thickness lmax melt, with lmax being the greatest thickness of the crystals that develop under the adopted crystallization conditions.

a lamellar morphology, which becomes largely imperfect with increasing hexene content. Imperfections typically correspond to the formation of distorted lamellae having small lateral dimensions, large distributions of the thicknesses of the crystalline and amorphous layers in the lamellar stacks, formation of short lamellar stacks, eventual inclusion of more than one population of lamellar stacks with different thicknesses, and, especially for the copolymers with higher concentration of hexene units, the presence of single lamellar entities besides a population of periodic arrays of parallel lamellae.29b The formation of distorted lamellar morphologies is typical of copolymers.35 In particular, the steep increase of the SAXS scattering intensity in the low q region of the SAXS profiles of the samples with hexene contents between 3.7 and 11.2 mol %, which are crystallized in the R form (curves c-e of Figure 4A), may be associated with the small lateral dimensions of the crystallites in these samples (Figure 3B), whereas the absence of this large scattering for the sample iPPHe18 in the trigonal form reflects the fact that for this sample the lateral dimensions of the crystallites become bigger because of the larger inclusion of hexene units in the crystals (Table 2 and Figure 3B). The average values of the lamellar long period L evaluated from the position q of the peak maxima in the Lorentz-corrected SAXS intensities of Figure 4A as L ) 2π/q were used for a rough evaluation of the thicknesses of crystalline layers lc through the product lc ) Lxc/100 with xc the crystallinity index determined from the wide-angle X-ray powder diffraction profiles of Figure 2.13,35 The values of long period L and thickness of crystalline layers lc are reported in Table 2 and in Figure 5A as a function of hexene concentration. It is apparent that with increasing hexene concentration the long period L first decreases from ∼118 Å for the sample iPPHe2.0 with 2.0 mol % of hexene to ∼105 Å


Figure 5. (A) Long period L (O) and crystalline lamellar thickness lc (b) of melt-crystallized compression-molded samples of iPPHe copolymers and of the homopolymer iPPA, determined from the SAXS profiles of Figure 4A, as a function of hexene concentration. (B) Experimental lc (b) and calculated lmax values of the crystalline lamellar thickness as a function of the end melting temperature (evaluated from the DSC curves of Figure 4B) of iPPHe copolymers with the indicated concentration (mol %) of hexene units (He). The experimental values lc are determined from the SAXS analysis (Figure 4A), and the theoretical values lmax are determined from the combined use of eqs 2 and 3, according to Crist’s method,40 using the end melting temperature Tm(lmax) values of Figure 4B and the following values for the model parameters: ∆Hm0 ) 209.5 J/g,41,42 Tm0 ) 186.1 °C,41c,d,42 σ ) 0.055 J/m,43 and F ) 0.95 g/cm3.44

for the sample iPPHe6.8 with 6.8 mol % of hexene, then increases, reaching the values of 135 and 180 Å for the samples iPPHe11.2 and iPPHe18, with 11.2 and 18 mol % of hexene, respectively (Figure 5A). The thickness of the crystalline layers lc, instead, undergoes a gradual decrease with increasing hexene content up to 11.2 mol %, from the value of ∼88 Å for the iPPA homopolymer to ∼42 Å for the sample iPPHe11.2, then increases with a further increase of hexene up to the value of 56 Å for the sample iPPHe18 with 18.0 mol % of hexene. These trends are similar to those outlined for the samples of iPPHe copolymers of ref 29 and reflect the different degree of inclusion of hexene units in the crystals. In fact, for hexene concentrations lower than 15 mol %, the samples crystallize in the R form and the hexene units are only partially included in the crystals,1,2,33 showing a decrease in the thickness and lateral dimensions of the crystallites as the hexene content increases. For concentrations higher than 15 mol %, a larger amount of hexene units are easily accommodated in the crystalline lattice of the trigonal form, producing a neat increase not only of the lateral dimensions of the crystallites but also of their thickness (Figure 5A and Table 2). The results of this analysis indicate that for iPPHe copolymers with hexene concentration lower than 18 mol %, which crystallize from the melt in the R form of iPP, the decrease of the melting temperatures (Figures 3A and 4B) is due to not only the increase of hexene concentration but also to the fact that, under the conditions adopted for the preparation of meltcrystallized compression-molded films, suitable for the analysis of the mechanical properties, the lamellar thickness also decreases (Table 2 and Figure 5A). However, the data of Figures


3A and 5A indicate that this is true only for the iPPHe copolymers with hexene content lower than 11.2 mol %, which crystallize in the R form. For these samples, indeed, a concomitant decrease of melting temperature and lamellar thickness with increasing hexene content occurs. For the sample iPPHe18 with 18 mol % of hexene, which crystallizes in the trigonal form, an increase of the lamellar thickness (Figure 5A), due to the high degree of inclusion of hexene units in the crystals, and a decrease of the melting temperature (Figure 3A) are, instead, observed. The melting temperature of this sample is lower than that of the sample iPPHe11.2 with 11.2 mol % of hexene that crystallizes in the R form because crystals of the trigonal form are intrinsically more defective than those of the R form owing to the higher degree of inclusion of hexene units in the trigonal crystals than that in the crystals of the R form.2 The experimental DSC and SAXS data of Figures 3A and 5A were analyzed on the basis of the theories of crystallization of statistical copolymers, in the effort to give an interpretation of the concomitant decrease of lamellar thickness and melting temperature with increasing hexene concentration. It is well known that the melting temperature of a copolymer crystallite depends on both its thickness and the content of counits.35b,36 As an example, it was shown in ref 37 that syndiotactic propylene/1-octene copolymers (sPPO) exhibit the same crystalline lamellar thickness as that of syndiotactic polypropene (sPP) homopolymer regardless of counit concentration, when they are isothermally crystallized from the melt at the same undercooling (with respect to the equilibrium melting point of sPP), even though, as expected, a shift of the melting point to lower temperatures is observed with increasing octene concentration. Since in sPPO copolymers octene units are excluded from crystals, this shift is well explained on the basis of Flory’s theory38 of copolymer crystallization, valid in the limit of strict exclusion. In fact, since the equilibrium melting temperature Tm is the ratio of the melting enthalpy ∆Hm to the melting entropy ∆Sm, Tm ) ∆Hm/∆Sm, for A/B random copolymers with dilute B units excluded from crystals of A units, the melting temperature of crystals is lower than that of the A homopolymer exhibiting the same crystal thickness because of different concentrations of the comonomeric units in the crystals in equilibrium with the melt and, therefore, to the presence of a nonzero mixing entropy contribution to the melting entropy. This contribution increases with increasing concentration of B content in the copolymer, producing a concomitant decrease of the melting temperature. It is worth mentioning that, according to the theory of Sanchez and Eby,39 the melting temperature of an A/B random copolymer would be lowered even in the case of inclusion of B units in the crystals. In fact, even in the limit of uniform inclusion of the B units in the crystalline and amorphous regions, which corresponds to a zero mixing entropy at melting, the enthalpy penalty for incorporating the B units in the crystals produces a decrease of the melting temperature. Within this context, it is useful to analyze the concomitant decrease of melting temperature and crystal thickness with increase of hexene content in iPPHe copolymers crystallized in the R form of iPP, using the method developed by Crist.40 The method has the advantage that it may be applied even to samples crystallized in conditions far from thermodynamic equilibrium, as those normally used for the determination of parameters of mechanical properties, does not involve any extrapolation procedure, and, most importantly, does not need any adjustable parameter.40 It allows establishing a lower (lmin) and an upper (lmax) bound for the crystalline lamellar thicknesses,

De Rosa et al.

which develop during crystallization in copolymers under the assumption of strict exclusion of defects from the crystals, using the results of simple DSC experiments (Figure 4B). Here the method was used to establish the values of lmax of our iPPHe samples to be compared with the values of the crystal thickness lc derived from SAXS measurements performed at room temperature, even though in these samples hexene units are partially included in the crystals. We demonstrate that Crist’s approach may be used as a test method to provide important structural information, even in cases of comonomer inclusion in the crystals. Crist’s method is based on the use of Sanchez and Eby theory of melting of crystals of statistical copolymers with finite crystallite thickness, containing an arbitrary concentration of comonomer units.39b According to this method, in the hypothesis that all the defects present in our samples, consisting of hexene units, stereoerrors rr, and secondary 2,1-erythro units (Table 1), are totally excluded from polypropylene crystals of R form, the end melting temperature observed in the DSC curves of Figure 4B (Tmax(lmax)) may be used to determine the largest thickness lmax of the crystals that develop in the crystallization process through the equation:

lmax )

(

Tmc 2σ F∆Hm0 Tmc - Tmax(lmax)

)

(2)

where F (in g/m3) is the crystalline density of hexene-free iPP crystals in the R form, and ∆Hm0 (in J/g) and σ (in J/m2) are the corresponding equilibrium melting enthalpy and the basal surface energy, respectively. Equation 2 can be rearranged to the familiar Gibbs-Thomson relation,36 in which Tmc corresponds to the equilibrium melting temperature of a copolymer with total molar fraction of defects xD, rather than to the equilibrium melting temperature of perfect (infinitely thick) iPP crystals of the R form Tm0. In other words, the values of Tmc to be used in eq 2 are evaluated using eq 3, which corresponds to the expression for the equilibrium melting temperature of copolymers derived by Flory38 in the limit of strict exclusion:

R ln(1 - xD) 1 1 ) 0c Tm Tm 42∆Hm0

(3)

where R is the gas constant, 42 is the molar mass of propene units, and xD is given by the sum of molar fractions of hexene units, stereoerrors rr, and secondary 2,1-erythro units (Table 1). The appropriateness of Crist’s approach is based on the following considerations. The melting point of copolymer crystallites depends on not only the concentration of counits and the lamellar thickness but also the composition of the coexisting melt. In the ideal case of complete exclusion of defects from the crystals of a random copolymer having the same thickness lc ) lmax, melting starts at a temperature well below Tm(lmax) due to a nonzero mixing entropy contribution to the melting entropy, because the concentration of defects in the melt is higher than the global concentration xD. As the melting proceeds, the concentration of defects in the melt gradually approaches the xD value and the last crystals melt at Tm(lmax). As a consequence, in the real cases, for copolymers approaching a random distribution of defects and total exclusion of these defects from the crystals, the values of lmax calculated using eq 2 in combination with eq 3 represent an upper bound to the average value of crystallite thickness lc that develops during crystallization and measured by SAXS at room temperature. In the case that defects are included in the crystals, instead,


Figure 6. Stress-strain curves of some compression-molded films of iPPHe copolymers with the indicated concentration of hexene (He) units. The stress-strain curves of the homopolymer sample iPPA is shown in enlarged stress and strain scales.

the values of lmax should be equal or lower than the values of lc measured at room temperature. Therefore, Crist’s approach40 may be used in a semiquantitative way as a test method to provide information concerning the inclusion/exclusion of the defects from the crystals, even in the cases of copolymer samples crystallized in non-isothermal conditions and far from equilibrium thermodynamics, as those generally used for measurements of the mechanical properties. For the calculation of the values of lmax of iPPHe copolymers, in agreement with the assumptions of the model that the last melting crystals consist of defect-free crystals of the homopolymer, the values ∆Hm0, Tm0, σ, and F correspond to those of R form of iPP and were assumed equal to 209.5 J/g,41,42 186.1 °C,41c,d,42 0.055 J/m,43 and 0.95 g/cm3,44 respectively. The so obtained calculated values of lmax are reported in Figure 5B as a function of the end melting temperatures of the meltcrystallized samples of iPPHe copolymers (Figure 4B), in comparison with the experimental crystalline lamellar thickness lc of the iPPHe samples determined from SAXS analysis (Figures 4A and 5A). It is apparent that only for the homopolymer sample iPPA the calculated value of lmax is higher than the lamellar thickness lc measured at room temperature, in agreement with the fact that the defects, essentially represented in this sample by secondary 2,1-erythro units,19a are largely excluded from the crystals of R form of iPP.19a For the copolymer sample iPPHe2.0 with 2.0 mol % of hexene, instead, the value of lc is nearly coincident with the theoretical value of lmax, whereas for iPPHe copolymers with hexene concentration higher than 2.0 mol % the experimental values of crystal thickness lc are slightly larger than the theoretical upper bound of the lamellar thickness lmax estimated in the hypothesis of full exclusion. This result is in agreement with previous experimental results that indicated partial inclusion of hexene units in the crystals of R form.2,33 Mechanical Properties. The stress-strain curves of compression-molded films of some samples of iPPHe copolymers are reported in Figure 6 and compared with the curve of the corresponding iPP homopolymer sample (iPPA) prepared with the same catalyst A. The values of mechanical parameters are reported in Table 2, whereas the values of Young’s modulus, deformation at break, and stress at yielding are plotted in Figure 7 as a function of hexene concentration. The highly isotactic homopolymer sample iPPA shows high value of the elastic modulus and very small deformability resulting in a highly stiff and fragile material.19a Incorporation of even small amounts of hexene units produces great increase


of deformation at break (Figures 6 and 7B) at values around 900-1000%, so that all samples of iPPHe copolymers show great enhancement of ductility, flexibility, and toughness compared to that of the homopolymer sample prepared with the same catalyst. The elastic modulus decreases with increasing hexene concentration (Figure 7A) according to the decrease of the degree of crystallinity (Figure 3B). Parallel to the decrease of the elastic modulus, the values of the stress at yielding strongly decrease with increasing hexene concentration (Figures 6 and 7C) and scales nearly linear with the crystalline lamellar thickness, at least for iPPHe samples with hexene content lower than 11.2 mol %, which crystallize in the R form (Figure 8). In particular, the more-crystalline samples with high crystalline lamellar thickness and low hexene content show plastic deformation via necking with a well-defined yielding, whereas lesscrystalline samples with low crystalline lamellar thickness and hexene concentrations higher than 9-10 mol % deform more homogeneously with ill-defined yield point. It is worth remarking that the linear relationship between stress at yielding and crystalline lamellar thickness of Figure 8 holds only for the homogeneous set of iPPHe samples that crystallize from the melt in the same crystalline form (i.e., up to hexene concentration of 11.2 mol %). In fact, the σy value of the sample iPPHe18 with 18.0 mol % of hexene that crystallizes in the trigonal form is lower than the σy value of iPPHe samples with lower hexene content, having the same crystalline lamellar thickness but crystallized in the R form. Since yielding corresponds to the disruption of the initial crystalline morphology, which sets in during the crystallization process,45 and is associated with the activation of a collective activity of crystallographic slip motions in the lamellae,46 the different behavior of the sample iPPHe18 reflects the fact that the intrinsic stability of defective crystals of the trigonal form of iPP including a high amount of hexene units is lower than the stability of the less defective crystals of R form, in which hexene units are less easily included. We recall that a linear relationship between yield stress and lamellar thickness has been found also in other polymers as for instance in the case of iPP homopolymer,47 polyethylene, and random copolymer of polyethylene,48,49 and several different mechanisms have been suggested to explain such behavior. The proportional relation of yield stress and lamellar thickness has been associated with the intrinsic complex molecular structure and morphology of semicrystalline polymers, and the most popular explanations include occurrence of partial melting and recrystallization,48 nucleation, and propagation of screw dislocations,50 the eventual role that thermally activated conformational (chain twist) defects play to overcome energy barriers for nucleation and propagation of dislocations.51 However, a definitive and general understanding of this proportionality and of the yield behavior of semicrystalline polymers in general has not yet been achieved.49 In a recent study, it was pointed out that, to reach a more definitive understanding, in addition to the structural and morphological aspects, kinetic factors associated with the segmental relaxation dynamics of the polymer chains should also be included in the theory.49 The data of Figure 8 suggest that an important parameter for a quantitative prediction of yield behavior of semicrystalline polymers is also the intrinsic stability of the crystalline modifications that develop upon crystallization. The mechanical properties of iPPHe copolymers after yielding (Figure 6) clearly indicate a decrease of resistance of crystals to the plastic deformation and easier deformability of samples with high hexene concentration. The values of tensile strength


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Figure 7. Average values of Young’s modulus (A), deformation at break (B), stress at yield point (C), and difference between stresses at break and at yielding (D) of iPPHe copolymers, evaluated from the stress-strain curves of Figure 6, as a function of hexene concentration. The values of modulus and deformation at break for the homopolymer sample iPPA are also reported.

Figure 8. Relationship between lamellar thickness lc and the yield stress σy of iPPHe copolymers with the indicated concentrations of hexene units (He). The copolymers with hexene concentration in the range 2-11.2 mol % are in the R form, whereas the copolymer with hexene concentration of 18 mol % (encircled) is in the trigonal form (Figure 2).

are instead nearly constant with hexene concentration (at least up to hexene content of 10-15 mol %; Figure 6 and Table 2). The high values of tensile strength, associated with low values of stress at yielding, are due to the strong strain hardening experienced by the samples at high deformation (Figure 6) and indicate that iPPHe copolymers show easy deformability but high strength. The difference between the values of stresses at break and at yielding, therefore, increases with increasing hexene concentration (Figure 7D). This parameter is a good indicator of the increased flexibility of iPPHe copolymers preserving high values of strength. Only at hexene contents higher than 15 mol %, when iPPHe copolymers crystallize in the trigonal form of iPP (samples iPPHe18 and iPPHe26),1,2 a strong decrease of elastic modulus and tensile strength is observed (Table 2). In these samples, the absence of the remarkable strain hardening during deformation, observed at lower hexene concentrations, accounts for the reduced values of tensile strength (Figure 6). These data indicate that iPPHe copolymers present mechanical properties of highly flexible thermoplastic materials with plastic resistance and modulus that decrease with increasing hexene content but high values of tensile strength. These copolymers allow tailoring stiffness and strength of the flexible material by

regulating the concentration of hexene comonomeric units. No elastic properties develop even in the case of poorly crystalline samples. This indicates that the properties depend on not only the comonomer concentration but also the type of comonomer (or, more generally, on the type of incorporated defect). For example, in propylene-ethylene copolymers, elastic properties develop for ethylene concentrations higher than nearly 10-12 mol %,52 and even in stereodefective iPP homopolymer samples the presence of rr stereodefects at concentration higher than 7 mol % induces development of elastic properties.16-18,20 In numerous studies performed to date on a large number of thermoplastic polymers at temperatures higher than the glass transition temperature, it has been shown that the strain hardening behavior, and in particular the slope of stress- strain curves at deformations close to the breaking, does not depend on the degree of crystallinity or lamellar thickness and appears primarily related to the density of chain entanglement and therefore to the intrinsic flexibility of the chains in the amorphous regions.48a,c,49,53 Morphological parameters may influence the strain-hardening behavior of these systems only indirectly.49 In particular, the higher the flexibility, the higher the density of chain entanglements and the strain hardening. For instance, it has been shown in the case of polyethylene that slowly melt crystallized samples exhibit a low strain hardening, associated with a low chain entanglement density, probably caused by a reeling in of the chains in a slow crystallization process.49 The nearly inactive role of crystals in the stress-strain behavior of semicrystalline polymers at deformations after yielding, after disruption of the crystalline phase and irreversible transformation of the spherulitic morphology into a fibrillar morphology, suggests that crystallites at high deformations do not act as physical cross-links, probably because they achieve a high mobility. Therefore, the large and nearly identical slope of the stress-strain curves in the strain hardening region experienced by iPPHe samples with hexene concentration in the range 2.5-11.2 mol % of Figure 6 may be somehow related to a high density of chain entanglement of these samples, which is probably higher than that of iPPHe samples with hexene concentration lower than 2.5 mol %. This behavior reflects the



Figure 9. X-ray fiber diffraction patterns of oriented fibers of copolymer samples iPPHe2.5 (A-C), iPPHe6.8 (D-F), and iPPHe11.2 (G-I) stretched at the indicated values of deformation ε. The (110)T reflection of the trigonal form is indicated in H and I.

fact that the intrinsic flexibility of iPPHe copolymer chains increases with increasing hexene concentration because the butyl side chains act as a sort of internal plasticizers. However, the absence of a remarkable strain hardening during deformation observed in the case of iPPHe samples with hexene concentrations higher than 11.2 mol %, which crystallize in the defective trigonal form, is an indication that the intrinsic stability of the crystals also plays a role in the ultimate mechanical properties of these samples. In fact, since in semicrystalline polymers the crystalline and amorphous phases may be viewed as two interpenetrating networks, the stress level transmitted to the crystalline skeleton by the entangled amorphous network may easily induce at high deformations various phenomena, such as phase transformations, local melting, or crystallization. These transitions are expected to occur more easily for crystals of low intrinsic stability. On the basis of these considerations, the mechanical properties of iPPHe copolymers can be largely explained on the basis of the structure, inclusion of hexene units in the crystals of R form of iPP,2 and structural transformations occurring during deformation. Inclusion of hexene units in crystals of R form2 produces large disturbance of the crystalline lattice and a consequent decrease of melting temperature, degree of crystallinity, crystallite thickness, and resistance of the defective crystals to the plastic deformation. This, in turn, induces a great increase of ductility compared to that of the homopolymer and a decrease of elastic modulus with increasing hexene concentration (Figure 7A). All samples show similar ductility, regardless of hexene concentration (Figure 7B), according to the similar values of crystallite size, at least up to nearly 4-5 mol % of hexene.

Moreover, the enhancement of flexibility and toughness compared to the highly isotactic homopolymer samples and the strong strain hardening that allows maintaining high values of strength may be also due to the occurrence of structural transformations during deformation. Stress-Induced Transitions. The X-ray fiber diffraction patterns of fibers of the samples iPPHe2.5, iPPHe6.8, and iPPHe11.2 with 2.5, 6.8, and 11.2 mol % of hexene stretched at various deformations, and keeping the fibers under tension, are reported in Figure 9. It is apparent that defective crystals of R form, present in the compression-molded films (Figure 2), transform by stretching into the mesomorphic form of iPP already at low values of deformation (100-200%), as indicated by the transformation of the three equatorial reflections (110)R, (040)R, and (130)R of R form into the broad halo in the range 2θ ) 14-16° in the diffraction patterns of Figure 9A,D. At high degrees of deformation, fibers in the pure mesomorphic form are obtained for iPPHe samples having hexene concentrations lower than 7-9 mol % (Figure 9C,F). We recall that the mesomorphic form of iPP is generally obtained by rapidly quenching the melt to very low temperatures and is characterized by chains in the ordered threefold helical conformation but high degree of disorder in the lateral packing of chains.54 The transformation of the R form into the mesomorphic form by stretching has also been observed in iPP homopolymer samples prepared with Ziegler-Natta55 as well as metallocene catalysts16,18 and has been suggested to occur through the destruction of the lamellar crystalline phase, probably by pulling chains out from crystals. Also for our metallocene-made iPPHe copolymers we can assume that the


formation of the mesomorphic form occurs via the pulling out of the chains from the lamellae of pre-existing crystalline form and successive reorganization of the chains in the crystalline mesomorphic aggregates. The dominant constituents of these mesomorphic aggregates are oriented bundles of chains, characterized by ordered threefold helical conformation but high degree of disorder in the lateral packing of chains.54 The easy formation of the mesomorphic form in iPPHe copolymers with hexene concentrations lower than 7-9 mol % already at low degrees of deformation (Figure 9A,D) is related to the incorporation of hexene units into the crystalline phase that introduces crystalline defects and structural disorder. These defective and small crystals of R form are less resistant to the plastic deformation and, therefore, can be more easily deformed and transformed into the mesomorphic form. Once the stiff crystals have been broken, the formation of the disordered mesomorphic form after yielding facilitates further stretching up to very high values of deformation giving high flexibility. The fast transformation into the mesomorphic form at low degrees of deformation and the absence of orientation of initial crystals of R form suggest that this transformation occurs through mechanical melting of original crystals and fast recrystallization into mesomorphic aggregates. The strain hardening observed in the stress-strain curves during deformation is associated with this polymorphic transformation and a possible increase of the degree of crystallinity, resulting in high values of the tensile strength. In the case of highly stereoregular homopolymer sample iPPA prepared with the same catalyst containing only negligible amount of rr stereodefects and 2,1 regiodefects, more ordered and big crystals of R form, present in the compression-molded films, have much higher plastic resistance and stability, so that most of the crystals cannot be easily deformed without breaking of the sample (Figure 6). This results in highly stiff and fragile materials.19a It is worth mentioning that stereodefective iPP homopolymer samples prepared with different C1-symmetric metallocene catalysts, containing higher concentration of rr defects (5-6 mol%),16 show similar flexibility and polymorphic transformations during deformation with the difference that compression-molded films are initially crystallized in the γ form that transforms by stretching into the mesomorphic form.16,18 A slightly different behavior was observed for iPPHe samples with hexene concentration higher than 7-9 mol %. In fact, the stretching of samples iPPHe9.0 and iPPHe11.2 produces already at low deformations formation of crystals of the trigonal form and of the mesomorphic form, even though the initial unoriented compression-molded films of these samples are crystallized in the R form (Figure 2h,i). The fiber diffraction patterns of the sample iPPHe11.2 are reported in Figure 9G-I as an example. It is apparent that during stretching the three equatorial reflections (110)R, (040)R, and (130)R of R form transform into the broad halo at 2θ ) 14-18° of the mesomorphic form, while the typical (110)T and (300)T reflections of the trigonal form at 2θ ≈ 10 and 17°, respectively, appear (Figure 9H). The intensity of the (110)T reflection of the trigonal form increases with increasing deformation, and at high deformation a fiber composed of a mixture of crystals of the mesomorphic form and the trigonal form (Figure 9I) is obtained. The broad halo at 2θ ) 14-18° is due to the superposition of the reflections of the mesomorphic form and of the (300)T reflections of the trigonal form. These data indicate that crystals of R form of the meltcrystallized films transform by stretching in part into the mesomorphic form and in part into the trigonal form. The reflections of the trigonal form are rather broad, indicating that

De Rosa et al.

small and quite disordered crystals of the trigonal form are obtained by stretching. Annealing of fibers of the sample iPPHe9.0 produces improvement of crystals of the trigonal form and, as expected, transformation of the mesomorphic form into R form of iPP.2 In the case of the sample iPPHe11.2, the mesomorphic form of iPP transforms by annealing completely into the trigonal form, rather than into the R form.2 This indicates that in stretched fibers of the sample iPPHe11.2 the trigonal form is more stable than the R form, even though this sample crystallizes from the melt only in the R form (Figure 2i). Stretching of iPPHe copolymer samples with higher hexene contents (iPPHe18 and iPPHe26), which are crystallized in the trigonal form (Figure 2l,m), only produces orientation of crystals of the trigonal form, and no polymorphic transformations occur during deformation. The incorporation into crystals of the trigonal form of very high amounts of hexene units and the corresponding formation of highly disordered crystals of this form and the absence of any polymorphic transformation during deformation account for the low strength of the samples iPPHe18 and iPPHe26 and absence of strain hardening. Conclusions The crystallization behavior, the stress-induced phase transformations, and the mechanical properties of highly isotactic propylene-hexene copolymers prepared with highly isoselective metallocene catalysts were analyzed and compared with the iPP homopolymer sample prepared with the same catalyst. Hexene units are partly included in crystals of R form of iPP at low hexene contents and largely included into the trigonal form at high concentrations. This produces large disturbance of the crystalline lattice and a consequent decrease of melting temperature, degree of crystallinity, crystallite size, lamellar thickness, and plastic resistance of the crystals. This, in turn, induces enhancement of ductility, flexibility, and toughness, compared to highly stereoregular iPP, and samples of iPPHe copolymers present mechanical properties of highly flexible material with high values of the tensile strength. The Young modulus and the stress at yielding decrease with increasing hexene content, according to the decrease of the crystallinity and of the crystalline lamellar thickness. A nearly linear relationship between the stress at yielding and the crystalline lamellar thickness, evaluated from the SAXS analysis, has been found. These outstanding properties are related to the structure and polymorphic transitions occurring during stretching. Defective crystals of R form, including hexene units, present in the compression-molded films show low resistance to the plastic deformation, which decreases with increasing hexene content, and rapidly transform by stretching at low degrees of deformation into the mesomorphic form of iPP. The formation of the mesomorphic form, in turn, facilitates successive further deformation of the sample up to very high strains with associated strong strain-hardening, resulting in highly flexible materials with remarkable high strength. Samples of iPPHe copolymers with high hexene contents, which are crystallized into the trigonal form, show low strength and absence of strain-hardening at high deformation. For these samples, stretching only produces orientation of crystals of the trigonal form. The incorporation into crystals of the trigonal form of very high amounts of hexene units and the corresponding formation of highly disordered crystals of this form and the absence of any polymorphic transformation during deformation account for the low strength and absence of strain hardening.


Acknowledgment. Financial support from Basell Polyolefins, Ferrara, Italy, is gratefully acknowledged.

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