Nucleation, structure and lamellar morphology of isotactic polypropylene filled with polypropylene-grafted multiwalled carbon nanotubes

European Polymer Journal 45 (2009) 2155–2163

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European Polymer Journal journal homepage: www.elsevier.com/locate/europolj

Macromolecular Nanotechnology

Nucleation, structure and lamellar morphology of isotactic polypropylene filled with polypropylene-grafted multiwalled carbon nanotubes Valerio Causin a,*, Bing-Xing Yang b, Carla Marega a, Suat Hong Goh b, Antonio Marigo a b

Department of Chemical Sciences, University of Padova, INSTM Research Unit, via Marzolo 1, 35131 Padova, Italy Department of Chemistry, National University of Singapore, 3 Science Drive 3, Singapore 117543, Singapore

a r t i c l e

i n f o

Article history: Received 27 November 2008 Received in revised form 21 May 2009 Accepted 21 May 2009 Available online 27 May 2009

Keywords: Nanocomposites Polypropylene Multiwalled carbon nanotubes Small and wide angle X-ray scattering

a b s t r a c t Polypropylene-based nanocomposites filled with polypropylene-grafted multiwalled carbon nanotubes (PP-g-MWNT) were compared to PP samples filled with pristine MWNT. The effect of such additives on the structure and morphology of the polymer matrix was studied by small angle X-ray scattering (SAXS), wide angle X-ray diffraction (WAXD), polarized light optical microscopy (PLOM) and differential scanning calorimetry (DSC). PP-g-MWNT allowed a more efficient and unhindered crystallization at a lamellar level, while MWNT disrupted the order of lamellar stacks, probably because of their tendency to aggregate. A common trend of tensile properties and lamellar morphology as a function of filler content was noted in the series filled with functionalized carbon nanotubes. Ó 2009 Elsevier Ltd. All rights reserved.

1. Introduction Since their discovery more than 15 years ago [1], carbon nanotubes (CNT) have attracted great interest in the academic and industrial research community, due to their outstanding electronic, electric, thermal and mechanical properties. Isotactic polypropylene (PP) is one of the most important commodity polymers and naturally drew a wide research activity aimed at the preparation of functional PPCNT nanocomposites. However, until now, reported property improvements have been less striking than expected [2–4]. Increases in modulus and strength are often coupled with decreases in ductility and toughness [5–9]. Just in a few cases the double aim of enhancing rigidity and ductility was accomplished. According to Zhao et al., the increase in ductility of their composites was due to increased mobility of PP chains and multi-wall carbon nanotubes (MWNT) due to orientation and to a bridging effect of oriented MWNT on the crack propagation during failure [10]. * Corresponding author. Tel.: +39 049 8275153; fax: +39 049 8275161. E-mail address: [email protected] (V. Causin). 0014-3057/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.eurpolymj.2009.05.026

Xiao and colleagues explained the simultaneous increase in tensile strength and impact strength with the uniform dispersion and possible orientation of nanotubes induced by the particular processing applied, i.e. dynamic packing injection molding [11]. Zhang and coworkers ascribed the remarkable properties of their samples to transcrystallinity [12]. Yang et al. attributed the improvement in tensile properties to the nucleation effect of CNT, to their uniform dispersion and to a more efficient transfer of stress due to the functionalization of MWNT [13]. To be effective reinforcing fillers, CNT must meet four main conditions: large aspect ratio, alignment, good dispersion and good stress transfer from the matrix to the filler, with the latter two being the most critical issues [2]. Bundling and aggregation of CNT have been considered the main hindrances for a homogeneous dispersion of the filler in the polymeric matrix. The importance of stress transfer between polymer and CNT is strictly related to the strength of interfacial interactions between these components of the composite. Polymer–CNT interactions may be chemical and van der Waals. Many approaches have been devised to enhance these interactions of CNT in the

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polymeric matrix, mainly based on either functionalization techniques or non-covalent wrapping methods. Chemical functionalization has the advantage of a permanent and very strong bonding of the functional groups on the CNT surface. However, it has the drawback of disrupting the conformation of the carbon atoms, therefore, decreasing CNT electrical and mechanical properties. The non-covalent method consists of wrapping CNT with surfactants, polymers or other molecules in order to modify their interactions with the matrix. Weaker links are obtained in this case, but the integrity of CNT is maintained. A number of papers have discussed the mechanism of reinforcement brought about by CNT [14–19]. Lordi and Yao suggested that the interfacial adhesion stems from a molecular-level entanglement of the two phases and forced long-range ordering of the polymer [14]. Coleman and coworkers stressed the importance of the crystalline phase nucleated by CNT in enhancing matrix-nanotube stress transfer [15,17,18]. Chang et al. advocated that formation of b phase in PP under strain had a role in reinforcing the composite [19]. Barber and coworkers reported that the properties of the polymer close to the CNT surface are different to those in the bulk [16]. Miltner observed that a transcrystalline phase of higher perfection than in the bulk formed around the nanotubes [20]. All these works point to the importance of studying the changes in structure and morphology of the polymer brought about by CNT, in order to understand the structure-property correlation governing the physical–mechanical behavior of these materials. Scanning and transmission electron microscopy (SEM and TEM, respectively), thermal analysis, Raman spectroscopy and wide angle X-ray diffraction (WAXD) are probably the most commonly employed techniques to carry out such investigations. In particular, DSC and WAXD are mostly applied for the investigation of the role of CNT on shaping the semicrystalline framework. It is reported though that the effect of nanofillers on the structure and morphology of nanocomposites is largely concentrated on the lamellar stacks [21–24], so small angle X-ray scattering SAXS should be the method of choice for the study of such aspects of these materials. SAXS allows the study of polymer lamellar morphology, which is very likely to change as a consequence of CNT addition and to influence very effectively the physical–mechanical properties of the composite. With the purpose of studying such effect, PP-based nanocomposites containing PP-grafted-MWNT (PP-gMWNT) [13] were considered. The influence of MWNT functionalization on the structure and morphology of PP were studied by SAXS, WAXD and differential scanning calorimetry (DSC). 2. Experimental 2.1. Materials Isotactic PP (Mn  50,000) from Sigma–Aldrich Co. Ltd., was used as the matrix. MWNTs (purity >95%, diameter within 10–20 nm, specific surface area 300 m2/g) produced by chemical vapor deposition was obtained from Shenzhen Nanotech Port Co. Ltd., China. Polypropylene-graft-maleic

anhydride (PP-MA, 0.6 wt.% of MA) with a melt index of 115 g/min (ASTM D1238, 190 °C, 2.16 kg) was obtained from Sigma–Aldrich Co. Ltd. and was used in the grafting procedure of MWNT. 2.2. Functionalization of MWNT and grafting of polymer on MWNT The details of the functionalization of MWNT and of the grafting of the polymer were previously reported by Yang et al. [13] The as-received MWNT were first treated with a 3:1 mixture of concentrated sulfuric and nitric acid. This mixture was sonicated in an ultrasonic bath and then diluted with cold distilled water. The acidified MWNT (MWNT–COOH) were separated out and washed repeatedly with distilled water. The acid content of MWNT– COOH was 0.09 mol.%, as determined using the method reported by Marshall et al. [25]. The sample was then allowed to react with ethylenediamine. PP-MA was grafted to MWNT–NH2 by melt blending in an Atlas Minimax mixer (LMM, USA). 2.3. Preparation of polymer composites PP-g-MWNT thus prepared were then blended with pristine PP to obtain composites of varying MWNT contents. The blending was also carried out in the Atlas Minimax mixer at a speed of 80 rpm at 175 °C for 30 min. Samples containing 1%, 1.5% and 2% (w/w) of grafted MWNT were prepared. For comparison, PP/MWNT composites of similar MWNT contents were also prepared under the same processing conditions using as-received MWNT. By thermogravimetry experiments, performed on a TA Instruments SDT 2960 apparatus, it was moreover checked that the onset of degradation was located beyond processing conditions. 2.4. Wide angle X-ray diffraction WAXD patterns were recorded in the diffraction angular range 10–50° 2h by a diffractometer GD 2000 (Ital Structures) working in a Seeman–Bohlin geometry and with a quartz crystal monochromator on the primary beam (CuKa1 radiation). The application of the least-squares fit procedure elaborated by Hindeleh and Johnson [26] gave the degree of crystallinity by weight which was then transformed in degree of crystallinity by volume (/WAXD) [27]. 2.5. Small angle X-ray scattering The SAXS patterns of the samples were recorded by an MBraun system, utilizing CuKa radiation from a Philips PW 1830 X-ray generator. The data were collected by a position sensitive detector, in the scattering angular range 0.1–5.0° 2h and they were successively corrected for blank scattering. A constant continuous background scattering was then subtracted [28], the obtained intensity values were smoothed in the tail region, and the Vonk’s desmearing procedure [29] was applied. The one-dimensional scattering function was obtained by the Lorentz correction

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2.6. SAXS data analysis The evaluation of the SAXS patterns according to some theoretical distribution models [30] was carried out referring to the Hosemann model [31], that assumes the presence of lamellar stacks having an infinite side dimension. 2.7. Differential scanning calorimetry (dsc) All the measurements were carried out with a TA Instruments mod. 2920 calorimeter operating under N2 atmosphere. Polymer samples weighing about 5 mg closed in aluminum pans were used throughout the experiments. Indium and tin of high purity were used for calibrating the DSC temperature and enthalpy scales. The T 0m of the samples was determined by the Hoffman and Weeks procedure [32]. The isothermal crystallizations were performed in the DSC by erasing the thermal history of the sample for 5 min at 200 °C and cooling down at the maximum controllable rate to the desired crystallization temperature. After a suitable time, the samples were heated again at 10 °C/min in order to obtain their melting temperatures. Crystallization temperatures were chosen so that the induction time before the beginning of crystallization was larger than the time requested by the apparatus to recover from the instrumental drift. The peak maximum of the melting endotherm was taken as the melting point. Application of the technique proposed by Marand and coworkers [33,34] confirmed the results obtained by the Hoffman and Weeks approach. The kinetics of crystallization was studied monitoring the heat evolved during the crystallization at each Tc. The Avrami equation [35–37] was used to correlate the fraction X of crystallized material with the time t:

X ¼ 1  exp½Kðt  t0 Þn 

ð1Þ

K is the kinetic constant of crystallization, n is a coefficient linked to the time dependence and the dimensions of growth of crystallites. The crystallization behavior of the polymer was also studied according to the relationship between chain folded crystal growth rate and undercooling proposed by Hoffman and Lauritzen: [38]

ln G þ

U Kg ¼ ln G0  RðT c  T 1 Þ T c DTf

ð2Þ

where G is the crystal growth rate, U* is a constant characteristic of the activation energy for reptative chain motion and is equal to 1500 cal/mol, R is the gas constant, T1 = Tg  30, Tg being the glass transition temperature, f is a correction factor equal to 2Tc/(Tm + Tc) and G0 is a pre-exponential factor. The constant Kg is given by: 0

Kg ¼

jT m b0 rre kDH

ð3Þ

where j is an integer equal to 2 in the second regime of crystallization and to 4 in the first and third regime of crystallization, b0 is the thickness of the surface layer, r and re are the interfacial free energies of the lateral and fold surface, respectively, k is the Boltzmann constant and DH is the enthalpy of fusion per mole of repeat units. 2.8. Polarized light optical microscopy (PLOM) The spherulitic morphology of the samples was studied with a Leica DM4000M polarized light microscope. The sample was placed between a microscope glass slide and a cover slip in an oven preset at 200 °C for 15 min, to ensure uniform melting and to delete their thermal history. Subsequently the samples were transferred to a Mettler FP82HT hot stage and isothermally crystallized at 135 °C. The photomicrographs were taken between cross-polarizers with a Leica DFC280 digital camera. 2.9. Measurement of tensile properties The tensile properties of the samples (dimensions 25  5  0.2 mm) were measured using an Instron Model 3345 mechanical tester at room temperature. The strain rate was 5 mm min1. Six measurements were performed for each sample. 3. Results and discussion 3.1. Tensile properties and filler dispersion Two series of samples were considered in this work. One was prepared using pristine MWNT, the other contained PP-g-MWNT. The tensile properties of these samples were already reported by some of us [13]. The samples with functionalized MWNT simultaneously increased their tensile strength, ultimate strain, toughness and Young’s modulus. On the other hand, the specimens prepared with pristine MWNT, displayed a more modest increase of tensile strength and modulus and a decrease of ultimate strain and toughness [13]. The effectiveness of reinforcement was evaluated according to Coleman [17] as the rate of increase of a particular tensile property with the volume fraction of filler. This value was much larger for the composites containing PP-g-MWNT than for those based on pristine MWNT. For the functionalized samples this value was 126 GPa, 5.6 GPa, 5.9 mm/mm and 0.59 GJ/m3, respectively for Young’s modulus, tensile strength, ultimate strain and toughness. The corresponding reinforcement values for the series containing pristine MWNT were 58 GPa, 1.3 GPa, 6.8 mm/mm and 0.11 GJ/m3 [13]. The larger values for modulus confirm that it is relatively easier to obtain improvements in stiffness, but it is hard to increase ductility, as reflected by the two opposite values of reinforcement for ultimate strain in the two series of samples. Yang et al. [13] explained the improvement in properties with a better dispersion of PP-g-MWNT, as shown by TEM observation, by the nucleation due to MWNT and by the better transfer of stress from the matrix to the filler. In Yang’s paper it can be seen by TEM and SEM that bundling and

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I1(s) = 4ps2 I(s), where I1(s) is the one-dimensional scattering function and I(s) the desmeared intensity function, being s = (2/k)sinh.

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aggregation could be avoided in PP-g-MWNT [13]. Individually dispersed PP-g-MWNT could be observed in the fracture surfaces of the composites containing up to 1.5 wt.% of PP-g-MWNT [13]. Beyond this threshold, small bundles of nanotubes began to appear, with a concurrent decrease of tensile properties [13]. Irrespectively of filler content, the composites prepared with pristine MWNT showed big bundles and aggregates [13]. One of the aims of this paper was to further study the effects of the fillers on the polymer matrix.

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3.2. Wide angle X-ray diffraction WAXD was firstly employed. Fig. 1 shows the diffractograms of blank PP matrix and of the two samples containing 1% filler. The patterns are very similar and have the typical reflections of a phase, confirming that CNT do not induce formation of c or b polymorphs [9,12,19,39–45]. Table 1 summarizes the results obtained by fitting experimental patterns using the procedure proposed by Hindeleh and Johnson [26]. The trend of WAXD crystallinity degree was quite similar for both sample series: addition of either MWNT or PP-g-MWNT slightly decreased the crystallinity of the matrix. Previous works reported conflicting results regarding the effects of CNT on the degree of crystallinity of PP. Most of the authors found, either by DSC or WAXD, that the crystallinity of the matrix remained unaltered [20,46–48], with a few exceptions that recorded an increase [44,49] or a decrease [42,48]. When a variation was found in the degree of crystallinity it was, however, quite modest. As will be shown later in this paper, CNT showed a nucleating effect that greatly increased the crystallization rate. This rapid crystallization did not allow the polymer to reach a highly regular semicrystalline framework and thus the composites showed a decreased crystallinity degree with respect to the pristine matrix. CNT are often reported to promote preferential orientation in the crystallization of a polymer matrix [49,50]. Interestingly, in our case addition of nanotubes of either type promoted the formation of more isotropic crystallites. Fig. 1 shows that in the diffractogram of the matrix, the (0 4 0) peak located at about 16.7 °2h is unusually more intense than the neighboring (1 1 0) and (1 3 0) peaks at about 14 and 18.5 °2h. This indicates a preferential orientation of the mole-

Fig. 1. WAXD patterns of samples blank PP, 1% PP-g-MWNT and 1% MWNT.

Table 1 Degree of crystallinity (/WAXD) of the considered samples. /WAXD (%) PP blank 1% PP-g-MWNT 1.5% PP-g-MWNT 2% PP-g-MWNT 1% MWNT 1.5% MWNT 2% MWNT

66 66 63 61 60 63 59

cules along the b crystallographic axis [49]. The intensity ratios of the peaks are the usual ones for the composite samples. Both pristine MWNT and PP-g-MWNT, therefore, promoted the formation of an unoriented crystalline phase. 3.3. Isothermal crystallization Previous non-isothermal crystallization measurements made by DSC by some of us [13], showed that during a cooling ramp from the melt at 10 °C/min, the peak crystallization temperature of the composites increased by at least 9 °C with respect to that of pristine PP. In order to better quantify the nucleating effect due to the addition of MWNT and PP-g-MWNT, isothermal crystallization studies were carried out.

Table 2 Equilibrium melting temperatures (T 0m ) and interfacial free energies of the fold surface (re) of the considered samples. Sample

T 0m (°C)

re (105 J/cm2)

PP blank 1% PP-g-MWNT 1.5% PP-g-MWNT 2% PP-g-MWNT 1% MWNT 1.5% MWNT 2% MWNT

214 199 197 199 193 191 192

2.3 0.91 0.91 0.88 0.96 0.89 0.89

Fig. 2. Hoffman–Weeks plot of sample 1% MWNT.

depicted in Fig. 2. The T 0m for the matrix polymer iPP falls within the 212–215 °C range reported in the literature [34]. The composites had lower T 0m than their corresponding base polymer. Also Zhou et al. [41] and Miltner et al. [20] reported similar decreases in the T 0m for PP-based nanocomposites with CNT. This is analogous with what was observed in the case of PP-clay nanocomposites [51,52], and reflects an intense influence of the filler on the crystallization of the matrix. The presence of CNT, and the nucleating effect associated to them (vide infra) increases the quantity of amorphous material, especially in the regions between lamellar stacks. It may be noted that the samples prepared with functionalized carbon nanotubes had a higher T 0m than those filled with pristine MWNT. This is indicative of a less disrupting role of PP-gMWNT, with respect to pristine MWNT, on the ordered and regular crystallization of the matrix. The kinetics was quantified by DSC according to the Avrami theory, obtaining plots like those of Fig. 3. The fitting was performed not considering the deviations from linearity due to the onset of secondary crystallization at longer times. The data obtained by the fitting of such lines are reported in Table 3. An n of about 2 was found for the pure matrix, indicative of 2-dimensional growth with simultaneous nucleation, in accord with literature data on neat PP [51]. The n’s of nanocomposite samples lied in an interval between 2 and 3, indicative of heterogeneous nucleation followed by spherulitic crystalline growth [53]. A trend was observed in composites towards increasing n. Nucleants are known to bring about an increase in n values [34,51,54]. The usual decrease in the kinetic constant k with decreasing DT, i.e. with increasing Tc, can be

It was preliminarily necessary to determine the equilibrium melting temperature (T 0m ) of the samples. This quantity corresponds to the melting temperature of a large stack of perfect extended chain crystals. T 0m is very important because, for an accurate study of crystallization kinetics, the supercooling DT, i.e. the difference between T 0m and the crystallization temperature Tc, must be known. The difference in free energy between melt and crystal phases depends in fact on DT, so only comparing polymers crystallized at the same DT one can be sure that they experienced the same driving force for crystallization. Table 2 summarizes the results obtained by fitting plots like that

Fig. 3. Avrami plot at selected temperatures of sample 1.5% PP-g-MWNT.

Table 3 Avrami parameters for the considered samples. PP blank

DT (°C)

n

ln K

86 83 80 77

1.8 2.0 2.1 1.9

3.2 6.3 8.1 9.1

1% PP-g-MWNT

1.5% PP-g-MWNT

2% PP-g-MWNT

1% MWNT

DT (°C)

n

ln K

DT (°C)

n

ln K

DT (°C)

n

ln K

DT (°C)

n

ln K

71

2.6

0.5

71

2.6

0.1

69

2.5

0.6 68

2.4

1.4

66

2.3

1.5 65

2.4

2.3

65

2.6

0.1

68 65

2.5 2.8

2.5 3.7 63

62

2.8

2.3

62

2.5

2.5

59 2.5

56 2.6

3.9

62

2.4

2.5

5.7

59

2.5

2.4

7.3

56

2.8

2% MWNT

DT (°C)

n

ln nK

63

2.3

0.3

60

2.3

1.4

57

2.5

3.0

54

2.6

5.7

51

2.7

8.4

48

2.6

10.5

DT (°C)

n

ln K

64

2.6

0.3

61

2.4

1.4

58

2.4

2.3

55

2.5

5.3

51

2.6

8.6

49

2.5

9.7

1.9

3.7

6.9

8.6 54

2.4

5.2

7.6 57

56

3.0

6.2 60

59

2.4

1.5% MWNT

6.8

9.1 53

2.6

9.2

53

2.6

8.6

50

2.5

10.0

50

2.6

10.9

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Fig. 5. Hoffman–Lauritzen plots of sample PP blank (d) and of the composites containing 1.5% MWNT (j) and 2% PP-g-MWNT (N).

Fig. 4. PLOM micrographs of samples (a) blank PP, (b) 1% PP-g-MWNT and (c) 1% MWNT isothermally crystallized at 135 °C. The metric bar corresponds to 100 lm.

observed in Table 3. Moreover, if different samples are compared, it can be seen that the matrix crystallizes more slowly than the other composites. These results are evidenced in Table 3. The very different kinetics did not allow to operate at the same undercooling for all the samples. If iPP were crystallized at an undercooling lower than 77 °C the rate would have been too slow for a feasible experiment, while on the other hand increasing the undercooling in the nanocomposites would have brought about too fast

a rate. This, per se, means that addition of CNT greatly enhances the crystallization rate, with respect to that of the pristine polymer, as amply reported [5,9,20,39,41–43,45– 50,53,55]. If the ln K of the nanocomposites are compared, two observations can be made. Within each series, the nucleating effect seems to be saturated [5,20,46,55] beyond 1.5% filler content, i.e. adding further filler does not bring about significant increases in the crystallization rate, probably because at that point the rate determining step becomes crystal growth rather than crystallite nucleation [39]. More interestingly, it can be seen that, at similar undercoolings, pristine MWNT are better nucleating agents than PP-g-MWNT, i.e. K is larger for the series prepared with pristine MWNT than with PP-g-MWNT. It therefore, appears that filler aggregates, which are present in the less homogeneous nanocomposites based on pristine MWNT, are more efficient in nucleating crystallization than the individual PP-g-MWNT. PLOM was used to further visualize the nucleating effect of the filler (Fig. 4). A spherulitic morphology was observed in all the samples, isothermally crystallized at 135 °C, but the spherulite size sharply decreases with the addition of PP-g-MWNT and even more of pristine MWNT. Spherulites are also more irregular in the nanocomposites, probably due to the rapid impingement related to the fast crystallization, even though some degree of transcrystallinity may also be present [12,20]. The interfacial free energy re, that may be estimated by the Hoffman–Lauritzen theory, is directly proportional to the work of chain folding, and can then be used as a measure for evaluating how easy it is for macromolecules to fold into lamellae. re was therefore useful for a mechanistic interpretation of the increase in the rate of crystallization. Fig. 5 shows the Hoffman–Lauritzen plots obtained for some considered samples. A significant break in the slope was observed only for sample 2% PP-g-MWNT, in the vicinity of the reported threshold undercooling for the II ? III regime transition (48 °C) [56]. All the other points pertaining to the remaining samples had an uniformly linear behavior and were thus assumed to crystallize according to regime III. re values (Table 2) could be calculated from Kg obtained by fitting the lines of Fig. 5, and using Eq. (3), substituting 4, 6.26 108 cm, 11.5 erg cm2 and 1.96 109 erg cm3, respectively for j, b0,

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Fig. 6. SAXS patterns (dotted lines) of (a) blank PP and PP-g-MWNT-based samples and of (b) pristine MWNT-based samples. Traces calculated during the fitting procedure (solid lines) are also shown.

Sample

CÅ

AÅ

DÅ

rC/C

rD/D

/SAXS (%)

r///SAXS

N

PP blank 1% PP-g-MWNT 1.5% PP-g-MWNT 2% PP-g-MWNT

143 182 187 175

60 55 59 57

203 237 246 232

0.29 0.31 0.34 0.37

0.22 0.25 0.27 0.28

71 77 76 76

0.09 0.10 0.07 0.07

30 5 5 5

r and DH [57]. re decreases significantly for the nanocomposites with respect to the base polymer. This means that the energy required to create a new nucleus is lower when CNT are present, and therefore, is the reason for the increase in crystallization rate [41,51,52 and references within]. 3.4. Small angle X-ray scattering The similarity of WAXD and DSC results for the two series of composites, compared to the starkly different tensile properties of these samples, suggests that some differences may be present in the morphology and structure at a different length scale. The simultaneous increase of strength, modulus, ductility and toughness in the PP-g-MWNT series should be ascribed not only to a better dispersion of the filler in the matrix, but could also be due to the influence of the filler on the lamellar framework of the matrix. Small angle Xray scattering contributed in elucidating this aspect. Fig. 6 shows the SAXS patterns obtained from the samples in both series. As may be seen, whereas for the sample prepared with PP-g-MWNT a neat SAXS signal was present, a featureless SAXS trace was produced by those filled with pristine MWNT. This is per se indicative of the detrimental role of unfunctionalized MWNT on the regular ordering at a lamellar level [24,43,45] and reflects an analogous effect brought about by lamellar nanofillers such as graphite or clay [21,23,52,58–60 and references therein]. The experimental traces due to the matrix and to the samples with PP-g-MWNT were fitted according to a method [30] which was shown [61,62] to reliably determine the thicknesses and distributions of the crystalline and amor-

phous layers, the long period and the crystallinity, along with their distribution, associated to lamellar stacks. The results are shown in Table 4. The model used for the best fitting was that of a variable finite stack with an asymmetric distribution. Addition of PP-g-MWNT increased the long period and the thickness of the crystalline layer and, consequently, the crystallinity of lamellar stacks. It can be seen that the crystallinities assessed by SAXS have larger values relative to those estimated by WAXD. This divergence can be explained considering the difference between the two techniques. SAXS is only sensitive to the crystalline regions organized in lamellar stacks, whereas WAXD allows the detection of all the regions contributing to the semicrystalline framework, including the amorphous phase located between the lamellar stacks. Therefore, WAXD crystallinity is lower because the contribution of crystalline domains is ‘‘diluted” by the interstack amorphous. The thickness of the amorphous layer did not change significantly. As a counterbalance to this perfectioning effect of PP-g-MWNT on the lamellar stacks, there is the decrease in the average number of layers in each stack, i.e. 30 in pristine PP and just 5 in the composites (Table 4). In other words, PP-gMWNT induce the formation of smaller stacks, but with a higher degree of crystallinity. This may explain why the crystallinity of the composites was found to decrease, with respect to the neat matrix, when measured by WAXD and to increase when evaluated by SAXS. This is consistent with a morphology of the composites in which very ordinate but small lamellar stacks are dispersed in a large interstack amorphous. The larger size of lamellar stacks in the pristine matrix and the absence of the nucleating effect due to CNT reduced the entity of interstack amorphous material, as reflected by the slighter difference between

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Table 4 Morphological parameters obtained by SAXS analysis of the samples. The thickness of the crystalline (C) and amorphous layer (A), the long period (D), and the crystallinity (/SAXS) along with their relative distributions (rc/C = rA/A, rD/D and r///SAXS) and number of lamellae (N) are shown.

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Fig. 7. Trend of the thickness of the crystalline layer of the lamellar stacks and of (a) Young’s modulus, (b) toughness, (c) ultimate strain and (d) tensile strength [13] as a function of filler content.

WAXD and SAXS crystallinity degree for this sample. The reduction in size of the lamellar stacks can be related, as in the case of smaller spherulites observed by PLOM, to the large increase in the rate of crystallization that does not allow the necessary time for the development of large structures. The effect of PP-g-MWNT on the PP lamellar morphology is important because with this filler the formation of an ordered lamellar structure was not hindered, but rather promoted. It must be said that the acid content of MWNT–COOH obtained was not very high (0.09 mol.%), and that the reaction between amino groups and maleated polypropylene is normally not complete, even though some degree of further reaction can have occurred when the PP-g-MWNT were blended with the PP matrix. Moreover, rather short PP chains were used for grafting, as evidenced by the high melt index of the employed MA-PP. All these factors do not allow to advocate a complete inclusion of grafted PP in the formation of the PP lamellar stacks. The particular morphology of such stacks, though, suggests that grafted PP has some role in shaping the semicrystalline framework of the material. Especially, the increase of the width of the distribution of the crystalline layer thickness and of the long period with increasing PP-g-MWNT content could be linked to the wider variability in the thicknesses due to a partial inclusion in the stacks of the PP moieties grafted on the nanotubes. Modulus was shown to depend on lamella thickness and crystallinity [63]. Fig. 7a shows that this correlation was valid in our samples as well. The trend of Young’s modulus as a function of filler content was the same as the thickness of the crystalline layers of the lamellar stacks. The two trends diverged for the sample containing 2% PP-g-MWNT because on one hand the thickness of the crystalline layer sharply decreased, whereas modulus lev-

eled off. This is likely an effect of bundling and aggregation of PP-g-MWNT, which has its onset for our samples at 2% filler content. Crystallinity seems to play a less important role in this case, since within the composite samples it remains constant, whereas significant variations in modulus are observed for increasing PP-g-MWNT contents. The simultaneous increase of tensile strength and elongation at break indicates the dominating effect of interfacial adhesion. Obviously many factors are critical for the tensile behavior of polymer-based nanocomposites, among which dispersion, matrix–filler interactions, spherulitic texture, micromechanical deformation processes, skin-core structures and morphological features [64]. It is not possible, based on the data gathered in this work, to single out all of these factors. It can be surely said that the degree of dispersion plays a significant role: all the measured tensile properties began to decrease in the composite containing 2 wt.% PP-g-MWNT, i.e. in the sample where insurgence of aggregation and bundling of the filler was noted. Some role of lamellar morphology in contributing to shaping the mechanical performance of the samples is suggested by Fig. 7, that shows that the trend of the thickness of the crystalline lamella as a function of filler content follows that of tensile strength, toughness and ultimate strain. An analogous common trend of tensile properties and of lamellar features was previously noted for polyethylene composites filled with polyethylene-grafted-MWNT [24].

4. Conclusions PP-g-MWNT were used to reinforce PP. The obtained nanocomposites exhibited not only improved stiffness and strength, but also increased ductility and toughness.

The grafting of PP on MWNT allowed for a more efficient dispersion of the filler in the matrix. The structure and morphology of composites based on plain MWNT and on PP-g-MWNT were characterized and compared by small angle X-ray scattering (SAXS), wide angle X-ray diffraction (WAXD) and differential scanning calorimetry (DSC). SAXS proved especially useful, because it showed that addition of PP-g-MWNT allowed a more efficient and unhindered crystallization at a lamellar level, while MWNT disrupted the order of lamellar stacks, probably because of their tendency to aggregate. These data show that filler dispersion, lamellar morphology and tensile properties should be considered interdependent and call for a deeper study of the connection of such factors. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25]

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