Metallocene Copolymers of Propene and 1-Hexene: TheInfluence of the Comonomer Content and Thermal Historyon the Structure and Mechanical Properties

J. M. LOPEZ-MAJADA,1 H. PALZA,2,3 J. L. GUEVARA,2–4 R. QUIJADA,2,3 M. C. MARTINEZ,5

R. BENAVENTE,1 J. M. PERENA,1 E. PEREZ,1 M. L. CERRADA1

1Instituto de Ciencia y Tecnologıa de Polımeros, Juan de la Cierva 3, 28006 Madrid, Spain

2Departamento de Ingenierıa Quımica, Facultad de Ciencias Fısicas y Matematicas, Universidad de Chile, Casilla 2777

3Centro para la Investigacion Multidisciplinaria Avanzada en Ciencias de los Materiales, Santiago, Chile

4Departamento de Quımica, Universidad Catolica del Norte, Avenida Angamos 0610, Antofagasta, Chile

5CTR Repsol-YPF, Mostoles, Madrid, Spain

ABSTRACT: The relationships between the structure and properties have been estab-lished for copolymers of propylene and 1-hexene synthesized with an isotactic metallo-cene catalyst system. The most important factor affecting the structure and propertiesof these copolymers is the comonomer content. The thermal treatment, that is, the rateof cooling from the melt, is also important. These factors affect the thermal properties,the degree of crystallinity, and therefore the structural parameters and the viscoelasticbehavior. A slow cooling from the melt favors the formation of the c phase instead ofthe a modification. Regarding the viscoelastic behavior, the b relaxation, associatedwith the glass-transition temperature, is shifted to lower temperatures and its inten-sity is increased as the 1-hexene content raises. The microhardness values are corre-lated with those of the storage modulus deduced from dynamic mechanical thermalanalysis curves, and good linear relations have been obtained between these param-eters.Keywords: a and c forms; isotactic polypropylene; metallocene catalysts; propylene–1-hexene copolymers; SAXS

INTRODUCTION

Isotactic polypropylene (iPP) is one of the thermo-plastic polymers primarily used in industry andhas bigger growth inside commodity polymers be-cause it presents interesting behavior: good envi-

ronmental resistance, the possibility of being proc-essed by different methods (extrusion, injection,and blow molding), and ease of recycling at a mod-erate cost. Moreover, iPP exhibits a lower densitythan other polymers such as high-density polyeth-ylene (HDPE; 5%), polystyrene (PS; 14%), and poly(ethylene terephthalate) (PET; 50%). This featureis really important for its extensive use and pro-duction all over the world because any object man-ufactured with iPP requires less material than thesame object made with some other polymers.1

Correspondence to: M. L. Cerrada (E-mail: mlcerrada@ictp.csic.es)

V

However, one of the inconveniences of iPP isits insufficient impact strength at low environ-mental temperatures near its glass-transitiontemperature. The glass-transition temperature ofiPP is about 0 8C, depending on variables such asthe crystallinity, crystal morphology, molecularweight, and level of isotacticity. To provide an en-hancement of the impact properties of iPP, itsblending with various types of elastomers hasgained much attention in recent years because ofits great industrial and commercial importance.Additionally, polypropylene (PP) has also beenblended with different polyethylenes (HDPE, low-density polyethylene, and linear low-density poly-ethylene). However, the problem with these blendsis the formation of incompatible phases and inter-phases. Consequently, the use of compatibilizingagents is necessary to attain an optimum blendand really improve the impact resistance. One pos-sibility for overcoming the difficulties of the misci-bility of the different polymer phases is reactorgranule technology, which leads to heterophasiccopolymers2 with an excellent stiffness/impactbalance3,4 and a wide range of ethylene contents(up to 25 wt %) as the rubber phase.

On the other hand, the polymerization andcopolymerization of polyolefins have undergonegreat advances with the discovery of metallocenecatalysts by Kaminsky and coworkers5–7 becauseof their great versatility. Concerning propenehomopolymerization, the use of these catalyticsystems has allowed the synthesis of PP withseveral tacticities (isotactic, atactic, syndiotactic,hemiisotactic, etc.), with only the structure of thecatalyst being changed.8,9 Moreover, these metal-locene catalysts allow the incorporation of rub-bery polyolefinic macromonomers, such as atacticpolypropene10 and poly(ethylene-co-propylene),11–13

into macromolecules, making feasible the coexis-tence of hard and soft segments within the chains.These novel copolymers constitute a very interest-ing approach to stiffness/impact resistance optimi-zation.

Another alternative for modifying the iPPstructure and, therefore, its properties withoutthe necessity of blending it with other polymersis the copolymerization of propene with a-olefinsof different chain lengths.14–16 The interest inthese copolymers is twofold: first, they might im-prove the final behavior of iPP, and second, thesenew materials obtained by this synthetic routemight exhibit interesting novel properties. In con-trast to the previously mentioned approach, thedevelopment of soft and hard domains, this meth-

odology primarily changes the iPP crystallinestructure, and so a good rigidity/toughness bal-ance and better transparency can be attained ata given optimum comonomer content.

The thermal treatment applied to polymers isimportant because the structure and mechanicalproperties strongly depend on the thermal historyimposed during processing. In particular, iPP cancrystallize into different crystalline structuresaccording to the crystallization conditions.17 Thedevelopment of the a, b, c, or smectic forms dependson the crystallization conditions18,19 and on thenumber of structural irregularities of the polymericchain.20

The most typical and most stable crystallinestructure, the monoclinic a form, was character-ized for the first time by Natta and Corradini.21

The hexagonal b modification appears only underspecial crystallization conditions or in the presenceof selective b nucleating agents. The orthorhombicc form has been found in the case of low-molecular-weight iPP17,22,23 and in random copolymers ofpropylene and a-olefins, favored when the propor-tion of the incorporated comonomer is increased.24

Moreover, the c modification is especially favoredin the case of iPP synthesized by metallocene cat-alysts because of the presence of errors homoge-neously distributed among the different polymerchains.25

The objective of this work was to study theeffects of the incorporation of 1-hexene and thethermal treatment on the final properties of pro-pylene–1-hexene copolymers synthesized by ametallocene catalyst. The structural and thermalcharacterization was carried out by X-ray diffrac-tion (profiles at wide and small angles) and calo-rimetric analyses, whereas the evaluation of themechanical behavior was performed by dynamicmechanical thermal analysis and microhardness(MH) determination.

EXPERIMENTAL

Polymerization

Two copolymers of propene and 1-hexene wereprepared with an isotactic metallocene catalystfor the direct copolymerization.15,26 The reactionswere carried out at 55 8C and 2 bar of pressure ofpropene, with a relation cocatalyst/catalyst (Al/Zr) of 1000, within a 1-L Buchi autoclave reactorfor 30 min, in a process-type slurry with toluene asthe solvent. For all the reactions, 5.0 � 10�6 mol

LOPEZ-MAJADA ET AL.

of the catalyst rac-Et(2-Me–Ind)2ZrCl2 was used,and methylaluminoxane was employed as the co-catalyst. The initial concentrations of the comono-mer in the feed were 0.13 and 0.32 mol/L, respec-tively, for the two copolymers, and the solventvolume was adjusted in such a way that the totalvolume inside the reactor was 500 mL. An iPPhomopolymer sample was prepared under thesame conditions for comparative reasons.

Characterization

The comonomer (1-hexene) content and the tac-ticity (% mm) of the samples, determined by 13CNMR spectroscopy, are shown in Table 1. 13C NMRspectra of the samples were recorded with a VarianInova 300 spectrometer operating at 75 MHz andat 90 8C. The samples were dissolved in o-dichloro-benzene, and benzene-d6 (20% v/v) was used forthe internal lock in a 5-mm NMR probe. The ex-perimental parameters were as follows: an acquisi-tion time of 1.5 s, a relaxation time of 4.0 s, and apulse angle of 748.

The molecular weights were determined by gelpermeation chromatography (GPC) in a WatersAlliance 2000 chromatograph equipped with adifferential refractive-index detector and a set ofthree columns.15 The columns were calibratedwith monodisperse PS standards. Analyses wereperformed at 135 8C with 1,2,4-trichlorobenzeneas the solvent at 1.0 mL/min. The weight-averagemolecular weights and molecular weight distri-butions are shown in Table 1.

Films were obtained via compression moldingin a Collin press between hot plates (160 8C forthe homopolymer and 140 and 120 8C for thecopolymers with low and high comonomer con-tents, respectively, to keep constant the differ-ence between their melting points and the tem-perature of the hot plates) at a pressure of 2 MPafor 4 min. Two different thermal treatments were

analyzed. The first thermal history, labeled S,consisted of a slow cooling (ca. 2 8C/min) from themolten state down to room temperature, at theinherent rate of the press, after the power wasswitched off. The second one, named Q, applied afast quench (ca. 60 8C/min) between plates refri-gerated with cold water after the melting of thematerial in the press.

Density determinations were performed at roomtemperature in a water–ethanol gradient columncalibrated with glass floats. The gradient of thedensities varied between 0.870 and 0.915 g/cm3.The degree of crystallinity obtained from densitymeasurements (fc

density) was calculated with thefollowing equation:27,28

f densityc ¼ ðqc=qÞðq� qaÞ=ðqc � qaÞ ð1Þ

Equation 1 is based on the assumption of a two-phase system, crystalline and amorphous, where qis the measured density of the sample and qa andqc are the densities of the completely amorphousand pure crystalline phases, respectively. The val-ues of qa ¼ 0.852 g/cm3 and qc ¼ 0.936 g/cm3 forthe amorphous and crystalline phase densities,respectively, were used.29,30

Wide-angle X-ray scattering (WAXS) patternswere recorded in the reflection mode, at roomtemperature, with a Philips diffractometer witha Geiger counter connected to a computer. Ni-fil-tered Cu Ka radiation was used. The diffractionscans were collected over a period of 20 min inthe 2h range of 3–438, with a sampling rate of1 Hz. The goniometer was calibrated with a sili-con standard. The X-ray determinations of thedegree of crystallinity were performed by sub-traction of the corresponding amorphous compo-nent31 in comparison with the totally amorphousprofile of an elastomeric PP sample.

The thermal transitions were also investigatedwith real-time X-ray diffraction experiments withsynchrotron radiation. These experiments wereperformed at the polymer beam line at Hasylab(Deutsches Elektronen Synchrotron (DESY), Ham-burg, Germany). The beam was monochromatized(wavelength ¼ 0.150 nm) by Bragg reflectionthrough a germanium single crystal. Two linearposition-sensitive detectors were used simultane-ously, one of them covering the approximate 2hrange of 10–308 and the other being set at asample–detector distance (in the direction of thebeam) of 235 cm, covering a spacing range of 5–55nm. Therefore, simultaneous WAXS and small-angle X-ray scattering (SAXS) data were collected.

Table 1. NMR and GPC Characterization

Sample1-Hexene(mol %)

Mw � 10�3

(g/mol)a Mw/Mnb Tacticity

PP1 0.0 63 2.0 83.0CPH2.1 2.1 69 2.1 85.8CPH5.5 5.5 67 2.0 85.0

a Weight-average molecular weight.b Weight-average molecular weight/number-average mo-

lecular weight.

COPOLYMERS OF PROPENE AND 1-HEXENE

Film samples of about 20 mg were coveredwith aluminum foil to ensure homogeneous heat-ing or cooling and were placed in the tempera-ture controller of the line in vacuo. A scanningrate of 8 8C/min was used. The scattering pat-terns were collected in time frames of 15 s, sothat we had a temperature resolution of 2 8C be-tween frames. The calibration of the spacings forthe detectors was performed as follows: the dif-fractions of a crystalline PET sample were used forthe WAXS detector, and the different orders of thelong spacing of a rat-tail cornea (65 nm) were usedfor the SAXS detector.

The thermal properties were analyzed in aPerkinElmer DSC-7 calorimeter connected to acooling system and calibrated with differentstandards at a heating rate of 20 8C min�1. Thesample weight ranged from 6 to 9 mg. The sam-ples were first heated from �50 to 180 8C at aheating rate of 20 8C min�1 and then cooled to�50 8C at the same rate. For crystallinity deter-minations, a value of 209 J/g was taken as theenthalpy of fusion of the perfect a modification ofiPP.28

The viscoelastic properties were measuredwith a Polymer Laboratories MK II dynamic me-chanical thermal analyzer working in a tensilemode. The temperature dependence of the stor-age modulus (E0), loss modulus (E@), and loss tan-gent (tan d) was measured at 3, 10, 30, and 50 Hzover a temperature range of �150 to 150 8C (thislast temperature was modified according to thecomonomer content) at a heating rate of 1.5 8Cmin�1. The specimens were rectangular strips2.2 mm wide, approximately 0.2 mm thick, andmore than 15 mm long. The apparent activationenergy values were calculated according to anArrhenius-type equation, with an accuracy of60.5 8C in the temperature assignment of tan d.The frequency dependence with the temperaturein the relaxation mechanisms associated withthe glass transition has also been considered tofollow an Arrhenius behavior, although it is dueto cooperative motions.32 This approximation canbe made without a significant error because therange of analyzed frequencies is low enough to befitted to a linear behavior such as that just men-tioned.

A Vickers indenter attached to a Leitz MH tes-ter was used to carry out microindentation meas-urements. Experiments were undertaken at roomtemperature (23 8C). A contact load of 0.98 N anda contact time of 25 s were employed. MH values

(MPa) were calculated according to the followingrelationship:33

MH ¼ 2 sin 68oðP=d2Þ ð2Þ

where P (N) is the contact load and d (mm) is thediagonal length of the projected indentation area.

RESULTS AND DISCUSSION

Structural and Thermal Characterization

Table 1 shows the main characteristics of thesynthesized copolymers. They have similar mo-lecular weights, polydispersities, and tacticities.Therefore, it will be considered in the followingthat these variables are not relevant for the dis-cussion of the changes in the different propertiesof the samples. On the contrary, the density val-ues, crystallinity, and thermal and mechanicalproperties depend considerably on the comono-mer content, as discussed later.

The degree of crystallinity obtained from den-sity measurements with eq 1 is shown in Table 2.The density and, therefore, crystallinity valuesdecrease with increasing 1-hexene in the copoly-mers. On the other hand, the slowly cooled sam-ples present higher density values than thosequenched. At first approximation, this featuremight indicate that a more perfect crystal struc-ture is attained in the former specimens.

Figure 1 shows the WAXS traces of iPP andtwo copolymers either quenched (top plot) orslowly cooled (bottom plot) from the melt. The Qsamples exhibit exclusively the characteristic dif-fractions of the a-monoclinic cell, whose diffrac-tion peaks appear at 2h values of 14.1, 16.9, 18.4,21.1, and 25.88 corresponding to the (110), (040),(130), (111), and (041, 131) reflections of the a

Table 2. Density, Overall Degree of CrystallinityObtained by Density (fc

density) and WAXS (fcWAXS),

Long Spacing (L), and Most Probable Crystal Size (lc)

Sample Density fcdensity fc

WAXSL

(nm)lc

(nm)

iPPQ 0.899 0.59 0.55 9.71 5.34CPH2.1Q 0.894 0.52 0.40 8.80 3.55CPH5.5Q 0.878 0.33 0.25 8.70 2.18iPPS 0.904 0.64 0.58 11.36 6.59CPH2.1S 0.895 0.54 0.50 10.75 5.38CPH5.5S 0.883 0.40 0.42 10.53 4.42

LOPEZ-MAJADA ET AL.

modification of iPP.34 The intensity of the differ-ent diffractions decreases as the comonomer con-tent increases in the copolymer, and consequently,the crystallinity exhibits the same tendency. Inaddition, diffractions are broadened and shiftedto lower angles; this indicates smaller and lessperfect crystallites as the comonomer content israised.

The diffraction profile corresponding to thequenched specimen of CPH5.5 shows rather dif-fuse diffraction peaks. In fact, only two widemaxima, centered around 15 and 218, can be ob-served. A quite similar profile can be observed forthe mesomorphic modification of iPP developed inefficiently quenched samples of this polymer.35,36

It seems, therefore, that specimen CPH5.5Q iscomposed mainly of mesomorphic entities, al-though a certain amount of coexisting monocliniccrystallites cannot be disregarded.

This intermediate-ordered mesomorphic struc-ture has also been found in copolymers of isotac-

tic propylene and 1-octadecene with a molar frac-tion of approximately 5 and higher14 under mod-erate quenching conditions. Therefore, it seemsthat the presence of the comonomer in the propenecopolymers increases the ability to produce themesomorphic modification because it is observedeven under rather mild quenching conditions.

The iPP subjected to a slow crystallizationprocess presents a well-defined peak at 19.98 cor-responding to the (117) reflection,37 which is char-acteristic of the c modification. However, this dif-fraction is quite subtle in CPH2.1-S and almostimperceptible in the CPH5.5-S copolymer. Thedetermination of the relative proportion of bothphases is not straightforward because the diffrac-tograms corresponding to the a and c forms ofiPP are rather similar. Therefore, several meth-ods have been proposed.25,38 A deconvolution pro-cedure, after subtraction of the amorphous com-ponent,23 has been used in this investigation. Forthat purpose, the amorphous halo of an elasto-meric PP sample31 has been used. The estimationof the X-ray crystallinity is now straightforwardbecause all the original diffractograms are nor-malized to the same total intensity. Subse-quently, the crystalline diffractograms are decon-voluted with Pearson VII profiles for the crystal-line diffractions. The proportions of the twomodifications are obtained from the relativeareas of the diffractions at 2h values of 18.5 and19.98. The results evidently undergo a slightchange with a different method or different pro-files for the fitting because of the diffraction over-lapping. Moreover, it is assumed with this proce-dure that, in the samples with 100% a or c modi-fication, the relative areas of the diffractions at18.5 and 19.98, respectively, in relation to thetotal crystalline area, are the same in both cases.With this procedure, from the total crystallinitydegree of 0.58 exhibited by sample iPPS, approxi-mately 85% corresponds to the c modification,and 15% corresponds to the a polymorph. Aquantitative determination has not been per-formed for the copolymers because the error thatcan be accomplished is quite significant.

It has been reported in the literature that thedevelopment of the c form is due to the interrup-tions in the PP isotactic sequences. These inter-ruptions enclose either regio- or stereodefects and,in the copolymers, the comonomer units. There-fore, when discontinuities are more frequent, thatis, at high comonomer contents, a high proportionof the c polymorph is developed.24,25,39 In this re-search, a conclusion involving the content of c

Figure 1. X-ray diffraction patterns, at room tempera-ture, of different samples for the two thermal treatments.

COPOLYMERS OF PROPENE AND 1-HEXENE

modification with the comonomer molar fractioncannot be attained, as previously mentioned. Thismight be due to differences in the protocol of speci-men preparation. Isothermal crystallization39 atdifferent temperatures and a much lower crystal-lization rate25 (6 8C/h) have been used in compar-ison with the slow cooling rate (2 8C/min) hereimposed.

A comparison of X-ray profiles of the quenchedspecimens and those slowly cooled points out thesmaller crystallinity developed in the former sam-ples, the differences being more important as thecomonomer content increases, as can be deducedfrom the values shown in Table 2. Moreover, thebetter resolved diffractograms suggest better or-dering and longer crystallites in the slowly cooledspecimens. These features have been confirmedby the estimation of the size of the ordered enti-ties by SAXS. The corresponding long spacingsare presented in Table 2, showing significantlyhigher values for the S specimens. From thesevalues and the total WAXS crystallinity of the sam-ple, the most probable crystallite size in the direc-tion normal to the lamellae can be estimated by theassumption of a simple two-phase model. The re-sults for the most probable crystallite size arealso listed in Table 2. The crystallites becomesmaller as the crystallization rate is raised dur-ing cooling because quenching limits the develop-ment and perfection of the crystalline entities.However, it is evident that the comonomer con-tent is the most important variable determiningthe crystalline morphology (crystallinity and crys-tal size), as expected, because the lateral branchesof 1-hexene are long enough to be excluded fromthe crystalline cell of iPP.39

Figure 2 illustrates the thermal behavior ofthe different samples during the first melting afterprocessing to analyze the crystalline structure thatdirectly affects the mechanical properties. On theone hand, a displacement to lower glass-transi-tion temperatures and melting temperatures isobserved with increasing comonomer content. Onthe other hand, a clear annealing peak is exhib-ited for the different samples, especially for thecopolymer with the highest comonomer content,CPH5.5. There is an important fraction of thesample that has crystallized during processinginto rather small and imperfect crystallites, pri-marily within the quenched specimens. There-fore, they are able to melt and recrystallize duringthe stay of samples at room temperature beforetheir analysis. These initially small crystallites areslightly enlarged, leading to the appearance of that

peak located at 40–50 8C. This feature is impor-tant on account of the influence of annealing onthe mechanical response27 because this stay atroom temperature for a few days leads to animprovement in the crystalline regions, increas-ing the crystal perfection and crystallinity, asobserved by a comparison of the degrees of crys-tallinity obtained from differential scanning calo-rimetry (DSC) during the first and second heat-ing processes (Table 3).

Moreover, the melting enthalpy is also signifi-cantly reduced when the 1-hexene content in-creases in the copolymer. The introduction of morecomonomer units hinders the chain regularityneeded for the crystallization process. Consequently,the crystallinity degree of the copolymers, esti-mated from the melting endotherm, is lowered asthe 1-hexene content increases (see Table 3).

The slowly cooled samples present glass-tran-sition temperatures lower than those found for

Figure 2. DSC first melting curves for the homopoly-mer and copolymers with both thermal histories at20 8C/min.

LOPEZ-MAJADA ET AL.

the Q samples (Table 3). This effect can be attrib-uted to the existence of two types of crystallites:those organized into an a lattice and those witha c arrangement, as confirmed by the X-ray re-sults. The c structure is composed of smaller andthinner crystals in comparison with those of thea modification. These thinner crystallites seem tolead to a lower movement restriction within theamorphous phase.

Another characteristic, shown in Figure 2, isthe splitting of the main melting peak found inthe slowly cooled samples, most likely due to thejust mentioned coexistence of the c and a crystal-lites and thus their corresponding melt processes.The melting that takes place at low temperaturesis related to the c phase, whereas the one associ-ated with the a modification occurs at a slightlyhigher temperature.39 The two melting tempera-tures have been estimated by deconvolution ofthe melting curve at that temperature interval,and their values are listed in Table 3. Althoughthe c crystallites are formed at higher tempera-tures, they are less stable than the monoclinicones and melt at lower temperatures.39

Figure 3 shows the second melting process foriPP and the two copolymers. The most importantfeature is the appearance of a cold crystallizationin CPH5.5 because this sample does not crystal-lize totally during the cooling process at 20 8C/min. It is also evident from this figure that thereis no annealing peak in the second melting of thethree samples, as expected. Moreover, the mainmelting endotherms do not present appreciablesplitting because at this relatively fast crystalli-zation rate (cooling at 20 8C/min), the proportionof the c phase formed is rather low (if any).

As commented previously, copolymer CPH5.5does not crystallize completely during the coolingat 20 8C/min. A more detailed analysis of thecrystallization behavior of this copolymer hasbeen performed via cooling from the melt at dif-

ferent rates. Figure 4 shows the melting curvesafter cooling from the melt at the indicated rates.First, the cold crystallization, which appearsaround 10 8C, diminishes in intensity as the cool-ing rate does in such a way that the specimencooled at 5 8C/min does not present appreciablecold crystallization. Moreover, the effect of theextent of crystallization on the glass-transitiontemperature is rather evident: it passes from�18 8C for the sample cooled at 40 8C/min to�10 8C for the one totally crystallized during thecooling process (at 5 8C/min). Additionally, at the

Table 3. Glass-Transition Temperature (Tg), Melting Temperature (Tm), and Degree of Crystallinity Obtainedby DSC (fc

DSC) for the First and Second Melting and Crystallization Temperature (Tc)a

Sample Tg,1 (8C) Tm,1 (8C) Tg,2 (8C) Tm,2 (8C) fc,1DSC fc,2

DSC Tc (8C)

iPPQ 1 122 �6 123 0.43 0.39 83CPH2.1Q �1 109 �8 108 0.34 0.28 65CPH5.5Q �9 80 �14 81 0.17 0.16 0iPPS �3 123, 127 — — 0.46 — —CPH2.1S �5 107, 114 — — 0.35 — —CPH5.5S �12 77, 90 — — 0.23 — —

a Subscripts 1 and 2 represent the first and second melting, respectively.

Figure 3. DSC second melting curves for the homo-polymer and copolymers crystallized at 20 8C/min

COPOLYMERS OF PROPENE AND 1-HEXENE

glass transition, the difference in the specific heatbetween the glassy and rubbery states increaseswith the crystallization rate (see Fig. 4), pointingout the progressive decrease in the crystallinity.

Synchrotron Measurements

In addition to the previous analysis of the longspacing for all the samples at room temperature(see results in Table 2), the slowly cooled speci-mens were also investigated by real-time variable-temperature experiments employing synchrotronradiation. Simultaneous SAXS/WAXS profiles wereacquired during the heating of the samples at 8 8C/min. The diffraction profiles corresponding to sam-ples iPPS and CPH5.5S are shown in Figures 5and 6, respectively.

Regarding the WAXS diffractograms, theyhave a poorer resolution than those presented inFigure 1, so now the determination of the rela-tive proportions of the a and c modifications iseven more difficult. It is clear, however, that inthe three S-type specimens, the c modification ispredominant.

The SAXS results appear to be more interest-ing. The values of the most probable long spac-

ing, deduced from the Lorentz-corrected SAXSprofiles, are shown in Figure 7 as a function oftemperature. Focusing our attention on sampleiPPS, we can observe three regions: an initialone, up to around 60 8C, in which the long spac-ing is almost constant; a second region, between60 and approximately 105 8C, with a moderateincrease in the long spacing; and a final one,with a very important thickening. This last re-gion coincides with the main melting endo-therm (see Fig. 2) and is ascribed to the usualmelting–recrystallization phenomena. On theother hand, the end of the first region (and thebeginning of the second) is approximately thetemperature at which the annealing peak isobserved.

In the case of sample CPH2.1S, the results arerather similar, with some minor differences: theinitial long spacings are somewhat lower thanthose for iPPS, and the final thickening regionoccurs at lower temperatures, corresponding tolower melting temperatures.

The behavior for sample CPH5.5S is, however,quite different because now the second and thirdregions mentioned previously seem to be coinci-dent, and the thickening is much smaller.

Nevertheless, the values of the long spacingare very similar for the three samples in the firsttwo regions. However, the crystal thickness,deduced from the long spacing and the overallcrystallinity, decreases considerably from iPP toCPH5.5 (see Table 2 for the room-temperaturevalues).

Additional information can be obtained fromthe analysis of the relative SAXS invariant.40–42

The temperature evolution of this invariant forthe three S-specimens is shown in the upper partof Figure 8. After an initial region of a very smallincrease in the invariant, it follows a more consid-erable increase, before the final sharp decreaseduring melting.

The lower part of Figure 8 shows the corre-sponding derivatives of the relative invariant.These derivatives are rather similar to the melt-ing endotherms (see Fig. 2), although the peaksare now smoother than the DSC curves. The rea-son may be either the much lower temperatureresolution of the synchrotron experiments (thediffraction profiles are the averages of every 15 s,the time frame) or a much smaller sensitivity ofthese SAXS results to the presence of a and ccrystals. Anyway, the overall temperature pro-files of the synchrotron and DSC results are re-markably similar.

Figure 4. DSC melting curves for the CPH5.5 copoly-mer after crystallization at different rates: 5, 10, 15, 20,and 40 8C/min (from top to bottom).

LOPEZ-MAJADA ET AL.

Dynamic Mechanical Properties

Figure 9 shows plots of E0, E@, and tan d as func-tions of temperature for the different specimens.The viscoelastic behavior of iPP and the copoly-mers is influenced by variables that affect the

crystalline structure,14 such as the comonomer

content, degree of crystallinity, thermal treat-ment, and, therefore, c-form content in the speci-

mens synthesized with metallocene catalysts.iPP obtained with Ziegler–Natta catalysts (iPP)

Figure 6. Real-time SAXS and WAXS profiles, obtained with synchrotron radiationfor CPH2.1S, for a melting experiment at 8 8C/min. Only one of every two frames isplotted for clarity.

Figure 5. Real-time SAXS and WAXS profiles, obtained with synchrotron radiationfor iPPS, for a melting experiment at 8 8C/min. Only one of every two frames is plottedfor clarity.

COPOLYMERS OF PROPENE AND 1-HEXENE

as well as metallocene PP presents three visco-elastic relaxations,32,43 as depicted in Figure 9.The c relaxation is due to the rotation of methylgroups, centered around �55 8C; the b process isassociated with the glass transition of the amor-phous regions, at about 0 8C; and the third relax-ation, named a, is related to movements involv-ing the crystalline phase.

The structural parameters (mainly the como-nomer content, crystallinity, and crystal size) doinfluence the location and intensity of the differ-ent relaxation processes. Looking at the tan dplot, we can clearly observe the a mechanism onlyin iPP for both treatments and in CPH2.1S. How-ever, it appears for all the different specimens ana-lyzed in the high-temperature side of the b processas a shoulder in E@ graphs and as a significantdrop in the E0 plots at temperatures higher thanthat associated with the glass transition. Fromthese plots and Table 4, a shift to lower tempera-tures is found in the location of the a process asthe comonomer content increases. The crystalli-tes become smaller and their number is dimin-ished with the 1-hexene content (Table 2), and

consequently, their motions can start at lowertemperatures.

On the other hand, the b relaxation (associatedwith the glass-transition temperature) is also de-pendent on the 1-hexene composition. The motionsin this case take place within the amorphous re-gions, whose content increases with the 1-hexenemolar fraction. Additionally, all the details relatedto improvements in crystallites (crystallinity, per-fection, and size) hinder this movement. Therefore,its intensity becomes significantly higher and itslocation is moved to lower temperatures as como-nomer incorporation is raised in the copolymers(see Fig. 9 and Table 4). Concerning the effect ofthe thermal treatment, it seems that the pres-ence of c crystallites in the slowly cooled speci-mens reduces the hindrance to motion within theamorphous regions, probably because of theirlower size in comparison with that presented bya crystallites. Thus, the location of this relaxa-tion is found at a lower temperature in the S

Figure 8. Temperature (T) dependence of the relativeSAXS invariants (top) and SAXS invariant derivative(bottom) for different S-specimens during the meltingexperiments.

Figure 7. Effect of the temperature (T) on the Lor-entz-corrected long spacing (L) for different S-speci-mens during the melting experiments.

LOPEZ-MAJADA ET AL.

samples at a given composition, similarly to theDSC results. In CPH5.5, this thermal historyinfluence is minimized because the Q specimen ispractically mesomorphic instead of being arrangedinto a crystallites.

The c relaxation of PP is observed for all thecopolymers. Its intensity is considerably reducedand its position is shifted to lower temperaturesas the 1-hexene content increases in the speci-mens (see Fig. 9 and Table 4). These featuresmight be related to a decrease in the steric effect

between the methylene groups pendant from thePP backbone responsible for this motion.

Mechanical Response

E0

Figure 9 also shows the effect of the comonomercontent for both thermal treatments on the elas-tic component of the complex modulus at 3 Hz. Aconsiderable decrease in E0 can be observed asthe 1-hexene content increases because of a re-

Figure 9. Temperature (T) dependence of the real and imaginary components of thecomplex modulus and tan d of different samples.

COPOLYMERS OF PROPENE AND 1-HEXENE

duction in the stiffness. This loss of rigidity ob-served in the copolymers can be ascribed to thediminishment of the crystallinity and crystallitethickness. This rigidity reduction is more impor-tant at high temperatures, above the glass tran-sition. The E0 values at room temperature arelisted in Table 4.

On the other hand, the effect of the thermaltreatment applied during processing is practi-cally negligible at subambient temperatures for agiven specimen. However, above the b relaxationassociated with cooperative motion within theamorphous regions, slowly cooled specimens ex-hibit slightly higher E0 values than samplesquenched from the melt, probably because of theirhigher overall crystallinity (see Table 5).

MH Measurements

MH is another significant mechanical magnitudein polymers, measuring the resistance of the ma-terial to plastic deformation, and accordingly pro-vides an idea about local strain. MH involves acomplex combination of properties (the elasticmodulus, yield strength, strain hardening, andtoughness). Its dependence on the comonomercontent and thermal treatment is quite analo-gous to that observed for E0 in the copolymersunder study, as depicted in Figure 10 and detailedin Table 5. Therefore, a direct relationship is com-monly found between the elastic modulus andMH,32 and the following empirical equation hasbeen proposed:

MH ¼ aEb ð3Þ

where a and b are constants and E is the elasticmodulus. This equation is also fulfilled by manysystems,27,44,45 from thermoplastic elastomers tovery rigid polymers, that is, in a very broadrange of MH and E values. Figure 11 shows a

good linear relationship between MH and E0 inthe different specimens analyzed for the twothermal treatments.

Impact Strength

Because of the practical importance of improvingthe impact resistance of iPP, the impact strengthof the studied samples has been estimated by theconsideration of such a magnitude related to thearea of the tan d curves. The correlation of theimpact resistance with dynamic mechanical be-havior has been indicated by many authors.3,46–49

The area under the tan d curve from �150 to 30 8Cprovides an estimation of the impact strength, al-though it is not a direct measurement. A linearcorrelation between the notched Izod impact strengthand the area under tan d has been observed for iPP/ethylene-vinyl acetate copolymer (EVA) blends,48

whereas an exponential relationship has beenfound for the iPP/EPDM blends.48 In the samplesstudied here, the incorporation of 1-hexene intoiPP by copolymerization leads to an increase insuch an area and, accordingly, enhances the insuf-ficient impact resistance of PP at a low tempera-ture, as shown in Table 5. Comparing these resultswith those shown by a commercial Ziegler–Natta

Figure 10. Variation of E0 and MH as a function ofthe comonomer content.

Table 4. Temperatures (Tan d Basis, 3 Hz)and Apparent Activation Energies (kJ/mol)of the Dynamic Mechanical Relaxations

Sample Tc Tb Ta Eac Ea

b Eaa

iPPQ �61.0 11.0 54.0 70 350 150CPH2.1Q �73.0 5.5 — 65 320 —CPH5.5Q �78.5 0.5 — 55 290 —iPPS �67.5 10.5 68.0 70 350 150CPH2.1S �78.0 1.0 46.0 60 310 —CPH5.5S �80.0 0.5 — 50 270 —

LOPEZ-MAJADA ET AL.

iPP49 and a commercial multiphasic PP,3 bothtoughened with a plastomer, we find that thearea under the tan d curve in these copolymersshows a larger increase than that observed in thecase of the aforementioned blends. The areas forthe iPP/plastomer and heterophasic iPP/plasto-mer, both with 25 wt % plastomer, are 4.5 and5.9, respectively, taking into account that thearea for the iPP homopolymer is 3.5 and that forheterophasic iPP is 4.0. Therefore, on the onehand, the metallocene iPP now analyzedpresents an enhanced impact strength with re-spect to that exhibited by the other two iPPs syn-thesized with Ziegler–Natta catalysts, probablybecause of the smaller crystallites that developedin the former specimen. On the other hand, theincorporation of the comonomer seems to be moreeffective than blending because the areas aremuch higher and the weight percentages are 4.1and 10.4 in CPH2.1 and CPH5.5, respectively,values considerably lower than the value of 25 wt %previously mentioned for the plastomer.

However, such an impact-strength improve-ment is accompanied by a considerably higher de-crease in rigidity and then in the modulus for thecopolymers in comparison with the blends. There-

fore, the choice of one polymeric material type oranother (copolymers or blends) has to be takenas a function of the best compromise between ri-gidity and toughness necessary to accomplish thefinal requirements for a specific application. Themetallocene block copolymers with soft and harddomains10–13 synthesized with rubbery polyolefinmacromonomers might also be a very interestingapproach.

CONCLUSIONS

From the structural, thermal, and mechanicalcharacterization of two copolymers of propyleneand 1-hexene (and of the corresponding homo-polymer) synthesized by an isotactic metallo-cene catalyst, the relationships between the struc-ture and properties have been established.

The most important factor affecting the struc-ture and properties of these copolymers is thecomonomer content. The thermal treatment, thatis, the rate of cooling from the melt, is also im-portant. These factors affect the thermal proper-ties, the degree of crystallinity, and therefore thestructural parameters and the viscoelastic behav-ior. A slow cooling from the melt favors the forma-tion of the c phase instead of the amodification.

On the other hand, the b relaxation, associ-ated with the glass-transition temperature, isshifted to lower temperatures, and its intensityis increased as the 1-hexene content raises. TheMH values are correlated with those of E0 de-duced from dynamic mechanical thermal analy-sis curves, and good linear relations have beenobtained between these parameters. The indirectestimation of the impact strength has shown thatarea values in the CPH copolymers are higherthan those observed in some reported blends

Figure 11. Relationship between MH and E0 in dif-ferent samples.

Table 5. E0 and MH at Room Temperature and Areaunder the Tan d Curve from �150 to 308C.

SampleE0

(MPa)MH

(MPa)

Area under theTan d Curve

(Arbitrary Units)

iPPQ 960 45.1 5.0CPH2.1Q 480 28.8 7.5CPH5.5Q 90 9.6 10.9iPPS 1300 50.0 4.8CPH2.1S 560 31.7 7.3CPH5.5S 110 9.7 11.7

COPOLYMERS OF PROPENE AND 1-HEXENE

with plastomers, although copolymerization leadsto a stronger drop in the modulus.

The authors are grateful for the financial support ofthe Comunidad Autonoma de Madrid, Ministerio deEducacion y Ciencia, CONICYT, and CONICYT/Con-sejo Superior de Investigaciones Cientıficas ExchangeCollaboration Program (project 07N/0093/2002, projectMAT2004-01547, project FONDAP 11980002 andscholarship to H. Palza, and project 2003CL0028,respectively). Support from the European Commission(COST Action D17, WG D17/0004/00) is also acknowl-edged. J. M. Lopez-Majada is grateful to Repsol-YPFfor its financial support. Additionally, the authorsthank G. B. Galland (Instituto de Quimica, Universi-dade Federal do Rio Grande do Sul, Porto Alegre, Bra-zil) for help with the NMR measurements and discus-sion. The synchrotron work (in the A2 polymer line ofHasylab at DESY, Hamburg, Germany) was supportedby the European Community Research InfrastructureAction under the FP6 ‘‘Structuring the EuropeanResearch Area’’ Program (through the IntegratedInfrastructure Initiative ‘‘Integrating Activity on Syn-chrotron and Free Electron Laser Science’’, contractRII3-CT-2004-506008). The authors thank the Hasylabpersonnel, especially S. Funari, for their collaboration.

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