Sessione 2 Struttura, proprietà e architettura molecolare

XVIII Convegno Italiano di Scienza e Tecnologia delle Macromolecole, Catania 16-20 settembre 2007

Sessione 2 Struttura, proprietà e architettura

molecolare

XVIII Convegno Italiano di Scienza e Tecnologia delle Macromolecole, Catania 16-20 settembre 2007 S2 O 1

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POLYMERIC FRAMEWORKS

A. R. Albunia, C. D’Aniello, C. Daniel, G. Milano, P. Rizzo, V. Venditto, G. Mensitieri,1

G. Guerra Università degli Studi di Salerno, Dipartimento di Chimica, Via Ponte don Mellilo, 84084 Fisciano (SA), Italy;

1Dipartimento di Ingegneria dei Materiali e della Produzione, Università di Napoli ”Federico II”, Piazzale Tecchio, 80125 Napoli, Italy e-mail: [email protected] The improvement of gas storage, recognition and separation techniques represents a strategic industrial and environmental objective. Some techniques are based on gas absorption on high surface amorphous materials, like e.g. activated carbons,1 crosslinked polymers2 or carbon nanotubes3. More selective techniques are based on inclusion of gas-molecules into cavities of crystalline materials leading to molecular complex crystalline phases (generally clathrate phases). The available nanoporous crystalline frameworks can present a large variety of chemical structures. Mostly studied are inorganic frameworks (e.g., zeolites)4-5 and, more recently, metal-organic frameworks,6-9 although the high molecular masses of the framework atoms can limit the gas weight uptake. In this respect organic frameworks, i.e. nanoporous organic crystals,10-13 often referred as “organic zeolites”, could be more promising. The δ crystalline phase of syndiotactic polystyrene (s-PS), 14,15 can be considered as a first case of polymeric framework. In fact, already at room temperature, this polymer crystal phase is able to include suitable gas-molecules leading to co-crystals. It is well known that s-PS is a polymer which forms co-crystals (both clathrate and intercalate) with several, mostly aromatic and/or halogenated, guest molecules. By suitable solvent extraction procedures, the guest molecules can be easily removed leading to the nanoporous δ form (monoclinic, space group P21/a; a = 1.74 nm; b = 1.18 nm; c = 0.77 nm; γ %=117°; density = 0.98 g/cm3)14 with a permanent cavity (of 120-160 Å3) per 4 monomer units (Figure 1).15 This low density crystalline framework rapidly absorbs volatile organic molecules, also if present in traces in air or water, eventually leading to host-guest molecular complex phases and hence is promising for applications in chemical separations 16-20 and molecular sensorics.21-22 In this communication, we show that gas molecules like for instance butadiene (boiling temperature at atmospheric pressure, Tb = − 5°C), carbon dioxide (Tb = − 78 °C) and ethylene (Tb = − 104 °C), also at room temperature, form real clathrate phases with s-PS. These clathrate phases have been characterized by X-ray diffraction and infrared linear dichroism measurements. Gas release kinetics are strongly reduced as a consequence of clathrate formation and can be further reduced by (010) orientation of the host phase, i.e., by placing the ac layers of close-packed enantiomorphous polymer helices nearly parallel to the film plane. The reported results indicate that this robust and easy-to-process polymeric crystalline framework is suitable for

gas storage, controlled release and for packaging requiring ethylene and carbon dioxide removal. a

b

L

RL

R

(002)

(010)

(210)-

½½

½

O =C =

O

0,1 0,1

0,1 0,1

Figure 1. Along c view of two adjacent unit cells of the host δ form of s-PS. The shown guest location and orientation into the cavity is based on X-ray diffraction and linear dichroism experiments. Suitable crystallization procedures can lead (002), (010) or (210) crystal planes nearly parallel the film surface. Guest release kinetics can be controlled by these uniplanar crystalline phase orientations. Recently it has also been shown that high porosity s-PS aerogels can be obtained by supercritical CO2 extraction of the solvent present in s-PS physical gels.23 The crystalline phase of the aerogels as well as their structure depend on the crystalline structure (co-crystalline or β) of the junction zones of the starting gel.23 When the polymer-rich phase of the gel is a co-crystal, the corresponding aerogel is formed by semicrystalline nanofibrils (fibril diameter range between 100-200 nm) presenting the nanoporous crystalline δ-phase while when the crystalline phase of the gel is the β phase, the corresponding aerogel is characterized by β-phase interconnected lamellar crystals (thickness in the range 200-400 nm).23 Chloroform-vapor sorption measurements at low chloroform pressure have shown that δ-form aerogels present the high sorption capacity characteristic of s-PS δ-form samples (due to the sorption of molecules as isolated guests of the host nanoporous crystalline phase) associated with the high sorption kinetics typical for


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areogels (due to the high porosity and hence high surface area).23 Thus, these new materials present a fast sorption kinetics while maintaining a good handiness. Sorption experiments have been conducted for s-PS aerogels with different crystalline forms and of different porosities. Nitrogen uptake experiments at 77K have confirmed that δ form aerogels present crystalline nanocavities as well as amorphous microporosity. The equilibrium uptake of organic molecules, from dilute aqueous solutions, for δ form aerogels is high (e.g., more than 5 wt% of 1,2-dichloroetane, DCE, from 1 ppm solutions) and independent of aerogel porosity. Gravimetric experiments, as well as FTIR conformational studies on absorbed DCE, clearly indicate that guest absorption from diluted solutions occurs only in the crystalline nanocavities. High-porosity δ form aerogels (P>98%) can also present substantial water adsorption from the amorphous microporosity, which is favourable to the pollutant absorption kinetics. In fact, apparent guest diffusivity increases of several orders of magnitude, with respect to the corresponding films, have been observed. References 1. Zhao, X.B., Xiao, B., Fletcher, A.J., Thomas, K.M.

J.Phys.Chem.B 2005, 109, 8880-8888. 2. Simpson, E.J., Koros, W.J., Schechter, R.S. Ind.Eng.Chem.Res. 1996, 35, 4635-4645. 3. Penza, M., Cassano, G., Aversa, P., Cusano, A., Cutolo, A., Giordano, M., Nicolais, L. Nanotechnology 2005, 16, 2536-2547. 4. Kuznicki, S.M., Bell, V.A., Nair, S., Hillhouse, H.W., Jacubinas, R.M., Braunbarth, C.M., Toby, B.H., Tsapatsis, M. Nature 2001, 412, 720-724. 5. Zecchina, A.. Bordiga, S., Vitillo, J.G., Ricchiardi, G.; Lamberti, C.; Spoto, G.; Bjorgen, M.; Lillerud, K. P. J.Am.Chem.Soc. 2005, 127, 6361-6366. 6. Eddaoudi, M., Li, H., Yaghi, O.M. J.Am.Chem.Soc. 2000, 122, 1391-1397.

7. Pan, L., Adams, K.M., Hernandez, H.E., Wang, X., Zheng, C., Hattori, Y., Kaneko, K. J.Am.Chem.Soc. 2003, 125, 3062-3067. 8. Kitaura, R., Seki, K., Akiyama, G., Kitagawa, S. Angew.Chem.Int.Ed. 2003, 42, 428-431. 9. Millward, A.R., Yaghi, O.M. J.Am.Chem.Soc. 2005, 127, 17998-17999. 10. Soldatov, D.V., Moudrakovski, I.L., Ripmeester, J.A. Angew.Chem.Int.Ed. 2004, 43, 6308-6311. 11. Blau, W.J., Fleming, A.J. Science 2004, 304, 1457-1458. 12. Atwood, J.L., Barbour, L.J., Jerga, A. Angew.Chem.Int.Ed. 2004, 43, 2948-2950. 13. Sozzani, P., Bracco, S., Comotti, A., Ferretti, L., Simonutti, R. Angew.Chem.Int.Ed. 2005, 44, 1816-1820. 14. De Rosa, C., Guerra, G., Petraccone, V., Pirozzi, B. Macromolecules 1997, 30, 4147-4152. 15. Milano, G., Venditto, V., Guerra, G., Cavallo, L., Ciambelli, P., Sannino, D. Chem. Mater. 2001, 13, 1506-1511. 16. Musto, P., Mensitieri, G., Cotugno, S., Guerra, G., Venditto, V. Macromolecules 2002, 35, 2296-2304. 17. Daniel, C., Alfano, D., Venditto, V., Cardea, S., Reverchon, E., Larobina, D., Mensitieri, G., Guerra, G. Adv. Mater. 2005, 17, 1515-1518. 18. Mensitieri, G., Venditto, V., Guerra, G. Sensors and Actuators B 2003, 92, 255-261. 19. Annunziata, L., Albunia, A.R., Venditto, V., Mensitieri, G.; Guerra, G. Macromolecules, 2006, 39, 9166-9170. 20. Buono, A., Immediata, I., Rizzo, P., Guerra, G. J.Am.Chem.Soc. in stampa 21. Giordano, M., Russo, M., Cusano, A., Cutolo, A., Mensitieri, G., Nicolais, L. Appl.Phys.Lett. 2004, 85, 5349-5351. 22. Giordano, M., Russo, M., Cusano, A., Mensitieri, G., Guerra G. Sensors and Actuators, B 2005, 109, 177-184. 23. Daniel, C.; Alfano, D.; Venditto, V.; Cardea, S.; Reverchon, E.; Larobina, D.; Mensitieri, G.; Guerra, G. Adv. Mater. 2005, 17, 1515-1518.


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CO-CRYSTALS OF SYNDIOTACTIC POLYSTYRENE WITH POLAR GUEST MOLECULES

C. Daniel, P. Rizzo, G. Guerra

Università degli Studi di Salerno, Dipartimento di Chimica, Via Ponte don Mellilo, 84084 Fisciano (SA), Italy; e-mail: [email protected] Introduction Syndiotactic polystyrene (sPS) is a stereoregular polymer presenting a very complex polymorphic behaviour (including four crystalline phases) and is capable of forming co-crystals (molecular-complex crystalline phases) with several low-molecular-mass guest complex. Guest molecules of sPS co-crystals present well defines three dimensional order with respect to the crystalline axes, even when volatile gases. Moreover, because sPS co-crystals depending on their preparation procedure can assume three different kinds of uniplanar orientation[1], it is possible to control the orientation of guest molecules not only in the microscopic crystalline phase but also in macroscopic films. On these bases, films presenting sPS/active guest co-crystals have been proposed as advanced materials, mainly for optical applications. This constitutes an innovative approach in the area of functional polymeric materials which are instead characterized by a disordered distribution of active groups into amorphous phases, by chemical bonding of active groups to the polymer backbone (functionalization), or by copolymerization of standard monomers with co-monomers containing suitable active groups. Although several dozen of the sPS co-crystals have been described, all the presented guests are apolar or poorly polar. In particular, the guest of highest polarity which has been presently reported in the literature is o-dichlorobenzene (µ = 2.5D). In this contribution it will be shown that by sorption in δ or in co-crystalline sPS samples of guests dissolved in suitable solvent carriers, clathrate phases can be easily obtained with molecules of high polarity having a volume lower than the upper limit observed for sPS clathrate phases (0.25-0.26 nm3)[2]. These co-crystalline materials presenting a stable and three-dimensionally ordered disposition of polar guests, characterized by also large first-order hyperpolarizability, could be in principle used for nonlinear optical or electrical applications. It will also be shown that a new host crystalline form (thereafter named ε) is able to form co-crystals with polar guest molecules with molecules having a volume larger than the sPS clathrate upper limit volume mentioned above [3]. Results The guest molecules considered in this contribution are listed in Table I. Polarized FTIR measurements have shown that exposure of uniaxially oriented δ-form films exposed at room temperature to the molecules listed in table I can lead to the formation of co-crystals only with nitrobenzene (NB)

and 4-nitroanisole (NA) while for all the other molecules co-crystals with sPS are not formed by simply exposure of δ-form samples. Table I: Molecular structure, molecular volume (Vm), and dipole moment (µ) values of guest molecules considered in this contribution.

Compound Molecular structure Vm (nm3) µ(D) nitrobenzene 0.170 4.0

4-nitro-anisole 0.208 4.6 4-(dimethyl-amino) benzaldehyde

0.231 5.1

trans-β-nitrostyrene 0.222 3.8 4-nitroaniline 0.161 6.2

trans-4-methoxy-β-nitrostyrene 0.250 4.6

4-(dimethyl-amino)-cinnamaldehyde 0.275 5.6

It has been observed that a different procedure, involving the sorption of guests dissolved in suitable solvent carriers (for example acetone), not only allows the formation of sPS co-crystals with NB and NA but also allows it with guest of high polarity such as 4-nitroaniline (µ =6.2 D), as well as for polar and bulkier molecules like trans-4-methoxy-β-nitrostyrene. Moreover polarized FTIR spectra show that for all the co-crystals, the guest molecules are oriented with their molecular plane nearly perpendicular to the crystalline polymer chain axis. For molecules with a volume larger than 0.25-0.26 nm3 such as 4-(dimethyl-amino)-cinnamaldehyde co-crystals have not been obtained. When sPS samples characterized by the new ε crystalline form are treated with solutions of the guests listed in table I in acetone, results show that sPS co-crystals can also be obtained even for a guest molecule of large volume such as 4-(dimethyl-amino)-cinnamaldehyde. Furthermore results show that with this host crystalline form, the guest molecular planes and dipole moments are oriented nearly parallel to the polymer chain-axes rather than perpendicular as observed for the δ host crystalline phase.

N(CH3)2 OH

O2N NH2

CHCHO(CH3)2N

CH3O NO2

NO2

CH=CHNO2

CH3O CH=CHNO2


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As, depending on the preparation procedure, it is possible to obtain an orientation of the host crystalline phase with polymer chain axes perpendicular to the surface of the film (orientation (002))[1b], a preferential of the guest molecular dipoles perpendicularly to the surface of the films can be achieved starting from the ε-form films (see Figure 1A).

Figure 1. Schematic presentation of the orientation of the co-crystal chain axes (c) and of 4-nitroaniline guest molecules, as obtained by NA absorption in ε-form (A) and δ-form (B) films, presenting an ideal perpendicular (002) orientation. This orientation makes these new materials suitable candidates for electrical poling processes, possibly leading not only to non linear optical but also to ferroelectric and piezoelectric properties.

Conclusions It has been shown that by suitable treatments of sPS samples characterized by the δ or the ε form, co-crystals can been obtained with guest molecules with a high dipole moment (up to 6.2D for 4-nitroaniline) and molecules presenting a non-zero first-order hyperpolarizability being relatively high (up to β = 30 x 10-30 esu for 4-(dimethyl-amino)-cinnamaldehyde). As a macroscopic orientation of these guest molecules can be achieved, films with these co-crystalline phases could present interesting properties for non-linear or electrical (ferroelectric and piezoelectric) applications. References 1. (a) P. Rizzo, M. Lamberti, A.R. Albunia, O. Ruiz de Ballesteros, G. Guerra, Macromolecules, 35, 5854 (2002) (b) P. Rizzo, A. Costabile, G. Guerra, Macromolecules, 37, 3071 (2004). (c) P. Rizzo, A. Spatola, A. De Girolamo Del Mauro, G. Guerra, Macromolecules, 38, 10089 (2005). 2. C. Daniel, N. Galdi, T. Montefusco, G. Guerra, Chem. Mater., in press. 3. P. Rizzo, C. Daniel, A. De Girolamo Del Mauro, G. Guerra, It. Pat. Appl. SA2006A22.

Filmthickness

Film plane

c

Fi lmth ickness

Film plane

c

δε

B A


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TEMPERATURE RESOLVED SAXS/WAXD OF PROPENE/1-PENTENE RANDOM COPOLYMERS

G.C. Alfonso1, F. Azzurri1, L. Boragno1,2, G. Costa2, M. Canetti3, P. Stagnaro2

1 Università di Genova, Dipartimento di Chimica e Chimica Industriale,Via Dodecaneso 31 - 16146, Genova 2 CNR – ISMAC Genova, Via De Marini 6 - 16149 Genova; e-mail: [email protected]

3 CNR – ISMAC Milano, Via Bassini 15 - 20133 Milano Introduction Contrary to the typical behavior of random copolymers, it has recently been found that, on increasing the comonomer content, C3/C6 copolymers exhibit, after an initial decrease of crystallinity, a substantial level of order up to compositions containing over 25 mol.% of hexene co-units.1-3 Inclusion of hexene co-units into the crystal lattice, destabilize the α-form but favor the formation of a new modification, which is isomorphous to form-I of it-poly(1-butene), as shown in Figure 1. Here we present clear evidences that also random copolymers with 1-pentene behave in the same way and crystallize in the new trigonal modification over a very wide range of compositions, at least between ca. 11 up to over 40 mol.% of pentene. Figure 1 – Model of the crystal structure of the trigonal form of i-PP. Experimental Room temperature and temperature-resolved 1-D SAXS/WAXD patterns were collected at beamline A2 of HASYLAB-DESY, Hamburg, on samples that were quenched from the molten state and aged at room temperature during at least one week. Samples were placed in a temperature controlled oven and were heated at 10°C/min. Diffraction data were collected in the 2θ range from 9.0 to 23.0° in the wide angle region and over the 0.02-0.1 nm-1 range of the scattering vector in the small angle domain. Sampling time was fixed at 15s, thus enabling us to collect diffraction patterns every 2.5°C. Results and discussion In Figure 2 the WAXD patterns collected at room temperature on copolymer samples at various contents of 1-pentene, which were crystallized from the melt and aged at room temperature for at least one week, are shown. It is evident that the position of the diffraction peaks does not correspond to that of the known α, β or γ-polymorphs of polypropylene. Three strong reflections at 2θ = ca. 10, 17.5 and 20.5° characterize the pattern in

the explored 2θ-range. According to the proposed trigonal cell, they correspond to Miller indices 110, 300 and 220+211, respectively. It is quite surprising that, notwithstanding the intrinsic disorder deriving from the random distribution of pentene co-units, the diffraction peaks are very sharp, particularly the one corresponding to the (110) plane. Figure 2 – WAXD patterns of C3/C5 random copolymers at various contents of 1-pentene co-units. A closer examination of the position of the diffraction maxima evidences that there is a small gradual shift of the peaks towards smaller angles on increasing the percentage of comonomeric units. This is directly related to the expansion of the lattice in the plane normal to molecular axis imposed by the need of hosting increasingly numerous units carrying the bulkier side groups. As shown in Figure 3, the extent of lattice expansion depends on the size of the substituent. At a given composition the values of a and b axes of the trigonal unit cell are larger for the copolymers with hexene than with pentene. In agreement with the hypothesis of inclusion of comonomeric units also the slope of the lines is smaller for the less bulky substituent. Figure 3 – a=b lattice constants of the trigonal cell. ○: C3/C6 copolymers;3 ●: C3/C5 copolymers.

8 12 16 20 24

0

500

1000

1500

2000

1823.631.936.440.7

C 5, m

o l .%

2θ, °

Inten

sity,

a.u.

300110

220+211

8 12 16 20 24

0

500

1000

1500

2000

1823.631.936.440.7

C 5, m

o l .%

2θ, °

Inten

sity,

a.u.

300110

220+211

1.5

1.6

1.7

1.8

1.9

0 10 20 30 40 50Comonomer, mol.%

a=b,

nm


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The temperature-resolved WAXD patterns of the copolymer containing 18 mol.% of 1-pentene shown in Figure 4 are consistent with the broad melting endotherms detected by DSC and provide further evidence of the low melting point of this polymorph. Upon heating, the intensity of all reflections begin to decrease as early as at around 40°C and the last traces of crystalline order disappear at ca. 60°C. It has been observed that all copolymers with 1-pentene content in the range between 11 and 40 mol.%, aged at room temperature during at least 10 days, melt in the same temperature range, with negligible dependence on composition. The thermal stability of the crystallites seems to be mainly dictated by the intrinsic structural disorder associated to the random nature of the copolymers rather than by their morphological features. In facts, very weak signals could be detected in the SAXS region for most of the investigated samples, thus suggesting that crystallites are separated by wide interphasic regions with density comparable to that of the disordered trigonal modification.

Figure 4 –Temperature-resolved WAXD of the C3/C5 copolymer containing ca. 18 mol.% of 1-pentene units. Heating rate: 10°C/min.

Conclusions Random C3/C5 copolymers crystallize in the low melting trigonale modification recently found in C3/C6 copolymers. According to expectations, the composition range in which this new polymorph develops begins at a comonomer content which is slightly higher than that found for the copolymers with hexene and ends at compositions exceeding, at least, 40 mol.% of 1-pentene. Acknowledgements Authors are grateful to the European Community for financing the Project II-05 047EC at Beamline A2, DESY, Hamburg and are indebted to Dr. S.S. Funari for his valuable assistance during experiments. References 1. C. De Rosa, F. Auriemma, P. Corradini, O. Tarallo, S. Dello Iacono, E. Ciaccia, L.Resconi J. Am. Chem. Soc. 2006; 128, 80-81 2. B. Lotz, J. Ruan, A. Thierry, G. C. Alfonso, A. Hiltner, E. Baer, E. Piorkowska, A. Galeski Macromolecules 2006; 39, 5777-5781 3. C. De Rosa, S. Dello Iacono, F. Auriemma, E. Ciaccia, L. Resconi, Macromolecules 2006, 39, 6098-6109

8 12 16 20 24

0

500

1000

1500

2000

37.552.5

67.582.5

T em p e

r a tu r e

, ° C

2θ, °

Inten

sity,

a.u.


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MULTI-TECHNIQUES APPROACHE FOR THE INVESTIGATION OF SHEAR INDUCED CRYSTALLIZATION

F. Azzurri, D. Cavallo, G. C. Alfonso

Università degli Studi di Genova, Dipartimento di Chimica e Chimica Industriale, Via Dodecaneso, 31 16146, Genova, Italy Introduction A good level understanding of Flow Induced Crystallization (FIC) has recently been achieved thanks to simple and ingenious experiments.1-3 The application of shear flow promotes the development of ordered nanostructures which are nucleation precursors and can survive for long times, even in the molten state, well above the melting point of the crystasl.4-6 In this framework, one of the main targets for scientists and technologists involved in the many facets of flow induced crystallization in polymeric materials is still represented by the full awareness of the coupled role of molecular characteristics and flow parameters on the solidification process. However, this relevant result can only be achieved by combining information acquired by means of different experimental techniques. Along these lines, we present preliminary results obtained by in-situ Rheo-optical experiments, in-situ Rheo-SAXS/WAXD measurements and Rheo-SALS analysis on a commercial Polypropylene grade submitted to medium intensity shear flow fields. Experimental The melt crystallization behaviour of isotactic polypropylene after the application of low-medium intensity shear flow fields has been investigated. The sample was a commercial grade product, PP T30G kindly provided by Montell (Italy). In situ Rheo-Optical experiments were performed by means of a Linkam CSS 450 Rotating Shearing Device equipped with a Polarized Light Optical Microscope and enabling a strict control of the applied thermo-mechanical history. The quantitative information was acquired by evaluating the nucleation density, N, obtained by counting the number of spherulitic growth centres per unit volume appearing after a given thermo-mechanical history. In order to quantify the neat effect of the flow field, the shear-induced nucleation density, Ns,

0s N-N=N eq.1 was evaluated by subtracting the nucleation density in quiescent conditions, N0, from that observed under shear. By setting the Linkam Shearing cell on the X-Ray Synchrotron radiation at Hasylab (Beamline A2, D.E.S.Y., Hamburg) Rheo-SAXS-WAXD data were collected; in-situ Rheo-SALS experiments were performed using a He-Ne laser light source and by collecting the depolarized scattering patterns by means of a CCD camera. In order to study the role of shearing variables on the subsequent crystallization process, a step shear of short duration was imposed on the molten polymer at a temperature, Ts, higher (or equal) than the observed melting point of the polymer. After cessation of flow, the

sample was quickly cooled (30°C/min) to a suitable crystallization temperature, Tc=135°C, at which the solidification took place isothermally in quiescent conditions. Results and discussion In Figure 1 the results obtained by performing in-situ Rheo-Optical experiments at different shearing temperatures are shown. As expected, for a given shear rate, the lower is shearing temperature the higher the shear induced nucleation density is. By applying very intense shear rate at relatively low shearing temperatures the formation of row nucleated morphology was observed. On the other hand, the existence of a shear rate-shearing temperature region in which the effect of applied flow is negligible can also be detected.

0

2

4

6

8

10

0.0019 0.002 0.0021 0.0022 0.0023 0.0024 0.0025

1/T, K-1

SR=3

SR=30SR=60

SR=100

R o w nu c l ea t ed mo r p ho l o gy

155°C2 00 °C

Ne g l i g i b l e e f f ec t s, N s<0 . 1N q

Figure 1 - Nucleation density as a function of reciprocal shearing temperature at different shear rates. Shearing time is 20 s. Through a well tested quantitative analysis of shear induced nucleation density data,7 the number of shear induced nuclei developed in the unit volume can be plotted as a function of shearing time for different temperatures and shear rates. The resultant linear plots allows one to quantify the rate of injection of shear induced nuclei, β, as defined in equation 2.

( ) ( ) ( ) ( ) ( ) sscqssscqcTot tγ,Tβ+TN=t,γ,TN+TN=TN && eq.2 Alike with other polymers,7 also in this case the rate of injection of shear induced nuclei scales with the shear rate according to a power law with exponent ≈2, independent from shearing temperature. 2a

0 γRTAexpMβ=β & eq.3

The existence of a shearing critical temperature, Ts*, at which, at any practically applicable shear rate, the flow effects are negligible has also been established. Rheo-SALS measurements quantitatively corroborate the data acquired by Rheo-Optical methods. The role of shearing time and shear rate on the overall crystallization


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kinetics has been investigated. In figure 2, the evolution of crystallinity as a function of time in samples sheared at 135°C using different shear rates is presented. As expected, an increase of shear rate produces a faster overall crystallization kinetics.

0

0.25

0.5

0.75

1

0 3000 6000 9000Time(sec)

Xrel

Quiescente

gdot3

gdot10

gdot20

gdot1

gdot0.1

gdot35

Figure 2 – Degree of crystallinity as a function of time at 135°C for sample sheared 3s at different shear rate. Experiments performed at a fixed shear rate demonstrate that the overall crystallization kinetics is faster when shear is imposed for longer times. Assuming that the integrated transmitted intensity is proportional to the relative degree of crystallinity and that Avrami equation for three-dimensional growth of pre-determined nuclei suitably describes the progress of crystallization, the time to reach 10% of relative crystallinity, t0.1, is related to the nucleation density by the equation:

30.1

-3tot tGK=N eq.4

Combining eq. 2 and eq.4, one has an indirect way of evaluating of the rate of injection of shear-induced nuclei, β. To disclose the earlier stages of flow induced nucleation and to integrate traditional approaches of detection of structural order in solidifying samples and of monitoring the development of crystallinity in quiescent conditions, also SAXS-WAXD data under constant shear flow and variable temperature regimes have been collected. Thanks to the coupled use of synchrotron radiation and the shearing device, reliable information on the degree of order, on the crystals’ size and orientation and on the overall flow induced crystallization kinetics have been acquired. As an example, the 2-D WAXD patterns collected during shearing at 30 s-1 and during the subsequent quiescent crystallization at 145°C are shown in Figure 3. It clearly appears that a certain level of local order and orientation is attained during the shear at 150°C, in the undercooled molten state. At the very beginning of solidification, the nascent crystals are characterized by a high degree of orientation. Subsequent development of order from the relaxing molten chains involves formation of randomly oriented lamellar crystals, thus causing a progressive decrease in the overall orientation. The analysis of preliminary SAXS experiments performed under severe shearing conditions demonstrate that application of flow fields promotes the formation of quasi-crystalline structures,

developing at temperatures as high as 170°C and able to survive in the molten state after the cessation of flow.

Figure 3 – Evolution of WAXD patterns during a shear induced crystallization experiment. Conclusions A good level of understanding of flow induced crystallization has been achieved by coupling different experimental techniques. Rheo-optical measurements performed on i-PP confirm results previously obtained on other polymers: the nucleation density increases on increasing shear rate and shearing temperature. The rate of injection of shear induced nuclei, obtained both from direct evaluation of nucleation density or from Rheo-SALS experiments, scales with shear rate according to a power law with exponent ≈2. Rheo-SALS experiments evidenced that the overall crystallization kinetics increases on increasing shear rate and by prolonging shearing time due to the effect of flow on the concentration of active nuclei. WAXD and SAXS patterns recorded during shearing and in the subsequent crystallization have demonstrated the presence of a certain level of order even when shear is applied in the molten state. Acknowledgements Authors are grateful to the European Community for financing the Project II-05 047EC at Beamline A2, DESY, Hamburg. The authors also acknowledge MIUR for supporting their research within a national research project PRIN 2004. References 1. Janeschitz-Kriegl, H.; Ratajski, E.; Steldbauer, M. Rheol. Acta, 42, 355-364 (2003) 2. Janeschitz-Kriegl, H. Macromolecules, 39, 4448-4454 (2006)

3. Somani, R.H.; Hsiao, B.; Nogales, A.; Srinivas, S.; Tsou, A.H.; Sics, I.; Baltà Calleja, F.J.; Ezquerra, T. Macromolecules, 33, 9385 (2000) 4. Somani, R.H.; Yang, L.; Hsiao, B.S. Phys. A., 304, 145-157, (2002) 5. Garcia Gutierrez, M.C.; Alfonso, G.C.; Riekel, C.; Azzurri, F. Macromolecules, 37, 478-485 (2004) 6. Azzurri, F.; Alfonso, G.C. Macromolecules, 38, 1723-1728 (2005) 7. Azzurri, F.; Alfonso, G.C.; Macromolecules, Submitted May 2007.


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A CLATHRATE FORM OF SYNDIOTACTIC POLY(P-METHYLSTYRENE) CONTAINING TWO DIFFERENT TYPES OF CAVITIES

O. Tarallo, G. Esposito, U. Passarelli, V. Petraccone Università degli Studi di Napoli “Federico II”, Dipartimento di Chimica, Via Cintia, 80126, Napoli, Italy; e-mail: [email protected] Introduction Syndiotactic poly(p-methylstyrene) (s-PPMS) co-crystallizes with several low molecular weight substances to form clathrates. The known structures can be divided in two different classes, α and β, both characterized by a s(2/1)2 helical conformation of the polymer chains with a repetition period of 7.8 Å. α class clathrates form centrosymmetric cavities, delimited by eight benzene rings belonging to two enantiomorphous adjacent helical chains. In their unit cells, the number of cavities is equal to the number of chains while the number of guest molecules per cavity can be one or two (see Figure 1A). In β class clathrates instead, the cavities in which guest molecules are hosted are delimited by four phenyl rings belonging to a chain and two phenyl rings belonging to an isomorphous chain related to the first one by a 21 screw axis. In these types of clathrates there is a double number of cavities in respect to α class clathrates and there is only one guest molecule per cavity. For this class, despite the same shape of the cavities, two different symmetries, monoclinic (see Figure 1B) or orthorhombic (see Figure 1C), have been found as far as the overall packing of the chains is concerned.

Figure 1. Schematic projections showing the packing of the chains and the placing of the cavities found in the crystal structures of the α (A) and β (B, C) class clathrate forms of s-PPMS. In the lower part of the figure only one couple of polymer chains delimiting the cavities are shown. R = right-handed, L = left-handed helices. In the present paper the crystal structure of the clathrate form of s-PPMS containing benzylchloride, in which in the same crystal are present both the two kinds of cavities described before, is presented. Results and discussion The X-ray fiber diffraction pattern of a uniaxially oriented clathrate sample of s-PPMS containing benzylchloride was indexed in terms of a monoclinic

unit cell (a = 26.2 Å, b = 11.8 Å, c = 7.9 Å and γ = 116°). The space group suggested was P21/a. Density data indicates in the unit cell a guest/polymer chain ratio of two. Following the hypothesis that this could be an α class clathrate containing two guest molecules per cavity, packing models have been found by molecular mechanic calculations. The best models obtained (almost isoenergetical) are reported in Figure 2.

Figure 2. Packing models proposed for the crystal structure of s-PPMS/benzylchloride molecular complex in the space group P21/a. In the c(b+a/2) projection only one couple of enantiomorphous polymer chains are shown. R = right-handed, L = left-handed helices. The shortest non-bonded distances between atoms are indicated in Å. Structure factors calculations showed, for each one of the models of Figure 2, that none of them gave a good agreement between observed and calculated data. When a statistical combination of the two models was done, a good agreement was obtained only for the equatorial layer line reflections. Successively, we considered the possibility of a disordered structure, in which polymer helices of a given chirality could be randomly substituted by others with opposite chirality. In this way, the equatorial projection of the structure would be kept unaltered (in order to save the good agreement obtained with the experimental data) while the layer lines structure factors would change. In fact, for the particular arrangement of the chains found in the models presented before it is possible to demonstrate that the substitution of a chain of a given chirality with an enantiomorphous one is feasible, provided that the methyl groups of the polymer helices are left nearly in the same positions. As far as the arrangement of the two guest molecules between the polymer chains along the b+a/2 direction in the unit cell is concerned, it must be noted that if we


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consider a couple of adjacent chains along the b+a/2 direction, the substitution of a chain of a certain chirality with an enantiomorphous one (keeping almost coincident the positions of the methyl groups of the helices) transforms the original α class cavities (delimited by two enantiomorphous helices) in a double number of β class ones (delimited by isomorphous polymer chains) (see Figure 3).

Figure 3. Projections perpendicular to the c(b+a/2) plane of a layer of chains in the crystal structure of the s-PPMS/benzylchloride molecular complex according the models of Figure 2. In (A) the …LRLRL… succession of chains delimiting the cavity that would be found for the ordered model is represented while in (B) a fault in the chirality of the chain in the middle is represented. α class cavities are schematically represented by the bigger gray ellipses while the smaller ones represent β class cavities. R = right-handed, L = left-handed helices. Consequently the couples of benzylchloride molecules related by an inversion center that were hosted in the original α class cavities must be “split” and arranged into new β class ones (being now related by a 21 axis). The feasibility of this new arrangement of the guests was then checked through molecular mechanics calculations considering the packing energy of an hypothetical clathrate form containing only β class cavities obtained by the arrangement of isomorphous helices according to the P21 space group in a unit cell with constants a =13.1 Å (a/2 of the experimental cell proposed before), b = 11.8 Å, c = 7.9 Å and γ =116°. The best models we obtained (see Figure 4 shown only for the case of left-handed helices) are characterized by the same energy of those reported in Figure 2 and the arrangements found for the guest molecules in their β class cavities do not differ significantly from the four found in the α class cavities of the models of Figure 2. Once the feasibility of the considered disorder was verified, we calculated the agreement between experimental and calculated structure factors considering a whole series of models characterized by a different degree of L/R disorder assuming, for simplicity, that the overall P21/a symmetry was kept.

The lowest discrepancy factor (R = 0.15) was obtained for statistical models in which about the 40% of molecules are hosted in β class cavities.

Figure 4. Packing models of minimum energy proposed for hypothetical ordered crystals of the s-PPMS/benzylchloride clathrate containing only β class cavities according to P21 symmetry. These models have been obtained minimizing the energy in a unit cell obtained cutting by half the a axis of the experimental unit cell. Along b+a/2 direction only one couple of isomorphous polymer chains are shown. L = left-handed helices. Analogous results are obtainable considering all right-handed helices. Conclusions The crystal structure of the clathrate form of s-PPMS containing benzylchloride is presented. This molecular complex represents, to our knowledge, the first case of a polymeric clathrate form in which in the same crystal are present two kinds of cavities having different shape and dimensions: α class cavities (containing two guest molecules) and β class cavities (containing one guest molecule). The contemporaneous presence of these two types of cavities is originated by a disorder in the arrangement of right- and left-handed helices in the crystals. The practicability of this disorder is related to the particular staggered layout of the layers of chains containing the cavities and to the fact that α class cavities are able to contain two guest molecules. References 1. V. Petraccone, D. La Camera, L. Caporaso, C. De Rosa, Macromolecules 33, 2610 (2000). 2. G. Esposito, O. Tarallo, V. Petraccone, Eur. Polym. J. 43, 1278 (2007). 3. V. Petraccone, D. La Camera, B. Pirozzi, P. Rizzo, C. De Rosa, Macromolecules 31, 5830 (1998). 4. D. La Camera, V. Petraccone, S. Artimagnella, O. Ruiz de Ballesteros, Macromolecules 34, 7762 (2001).


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DETERMINAZIONE DELLA STRUTTURA CRISTALLINA DEL CIS-1,4-POLI(1,3-ESADIENE) ISOTATTICO MEDIANTE DIFFRAZIONE AI RAGGI X E MECCANICA

MOLECOLARE

B. Pirozzi1, R. Napolitano1, G. Giusto1, G. Ricci2 1 Università degli Studi di Napoli Federico II, Dipartimento di Chimica “Paolo Corradini”, Via Cintia, 80126, Napoli, Italy; e-mail: [email protected]

2 Istituto per lo Studio delle Macromolecole-C.N.R., via E. Bassini 15, 20133, Milano, Italy Introduzione Gli 1,3-dieni, possedendo due doppi legami, possono addizionarsi con concatenamento 1,4 o con concatenamento 1,2; si possono così ottenere polimeri che si differenziano per la posizione del doppio legame. Nel caso dell’addizione 1,4 si ottengono polidieni nella cui catena principale sono presenti doppi legami, aventi configurazione cis o trans, seguiti da tre legami semplici. Di conseguenza, i polidieni aventi configurazione regolare vengono classificati come cis-tattici o trans-tattici. Altri potenziali centri di stereoisomeria sono i due atomi di carbonio Csp3 della catena principale, per cui i polidieni configurazionalmente regolari possono essere (di)isotattici e (di)sindiotattici. In questo lavoro viene riportata la struttura cristallina del cis-l,4-poli(1,3-esadiene) isotattico [icisPED14] determinata mediante l’uso combinato della meccanica molecolare e della diffrazione ai raggi X. Risultati La sintesi dell’icisPED14 è stata condotta in eptano a -30°C partendo dall’isomero E dell’1,3-esadiene. È stato utilizzato un catalizzatore a base di neodimio e alluminio alchili. Questo sistema catalitico, associato alla bassa temperatura di polimerizzazione, ha permesso la sintesi di un polimero di alto peso molecolare e altamente stereoregolare. La sua densità, determinata per flottazione, è risultata 0.88 g/cm3. Il DSC, effettuato alla velocità di 10 °C/min, mostra la fusione del campione a 91°C, nessuna cristallizzazione durante il raffreddamento fino a 0°C e la cristallizzazione a 8°C durante il successivo riscaldamento. Il campione tal quale è stato sottoposto ai seguenti trattamenti: a) fusioni sotto pressione seguite da raffreddamenti all’aria a varie velocità; b) fusioni seguite da cristallizzazioni isoterme a varie temperature. Fibre sono state ottenute per stiri a varie temperature di campioni raffreddati dal fuso non isotermicamente. Successivamente le fibre sono state anche ricotte a varie temperature per tempi diversi. I diversi campioni ottenuti sono stati analizzati mediante la diffrazione ai raggi X. Dall’analisi dello spettro di fibra sono stati ricavati il valore dell’asse di catena c che è risultato pari a (8.10±0.05)Å e gli altri parametri (a, b e γ) di possibili celle ortorombiche e monocline. Sono stati effettuati calcoli di meccanica molecolare utilizzando il campo di forza Compass1. Per tutti i calcoli di meccanica molecolare, come per le simulazioni degli spettri di diffrazione ai raggi X, è stato utilizzato il programma Cerius2 di Accelrys. L’energia conformazionale è stata calcolata sia come mappa di una molecola modello che mediante minimizzazioni dell’energia di una catena polimerica avente conformazione regolare. La molecola

modello scelta è il 4-etil-(Z,Z)-2,6-ottadiene, contenente tutte le interazioni intramolecolari significative di una catena dell’icisPED14. I punti della mappa sono stati ottenuti variando ad intervalli di 2° gli angoli di torsione in catena θ1 e θ3 e minimizzando per ogni punto tutti gli altri parametri conformazionali. La mappa dell’energia conformazionale in funzione di θ1 e θ3 è riportata in Figura 1 insieme ad uno schema della molecola modello. Essa presenta due minimi, il più profondo dei quali corrisponde ad una conformazione avente parametri interni simili a quelli di una catena polimerica con conformazione elicoidale s(2/1).

Figura 1. Mappa dell’energia conformazionale della molecola modello schematizzata, in funzione di θ1 e θ3. Le energie, in kJ/mol, sono riferite al minimo della mappa assunto come zero. La minimizzazione dell’energia della catena polimerica isolata sotto il vincolo dell’asse c sperimentale ha fornito i valori di tutti i suoi parametri interni. I valori degli angoli di torsione più significativi sono riportati in Tabella I in confronto con quelli ad essi corrispondenti nel minimo della mappa della molecola modello. Tabella I

θ1 θ2 θ3 θ4 catena

polimerica -114.5° 175.1° 120.4° -63.4° molecola modello -118.0° 173.1° 116.0° -63.9°

Per completare la determinazione della struttura cristallina è stato assunto che l’elemento di simmetria 21

0 60 120 180 240 300 3600

60

120

180

240

300

360

θ1 (deg)

θ3 (de

g)

1.8

0

>10


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della catena sia anche cristallografico, per cui sono stati analizzati i gruppi spaziali monoclini ed ortorombici aventi come elemento di simmetria l’asse 21 non combinato con altri elementi di simmetria. Sono stati considerati soltanto i gruppi spaziali che prevedono la presenza di due catene nella cella elementare. Sono state effettuate minimizzazioni dell’energia in funzione sia dei parametri interni delle catene che delle loro posizioni relative, sotto il vincolo dei parametri delle possibili celle elementari ricavate sperimentalmente. Per ciascuna struttura ottenuta dalle minimizzazioni sono stati calcolati gli spettri di polveri e di fibra che sono stati confrontati con quelli sperimentali. Il migliore risultato, sia per l’energia sia per il confronto tra gli spettri sperimentali e calcolati, si è ottenuto con una struttura cristallina caratterizzata da una cella elementare ortorombica con assi a=9.0Å, b=7.8Å, c=8.10Å, nel gruppo spaziale P212121. Il confronto tra lo spettro calcolato per la struttura proposta e gli spettri di polveri sperimentali dei vari campioni è riportato in Figura 2. Figura 2. Spettro di polveri calcolato per la struttura cristallina proposta (a) in confronto con gli spettri delle fasi cristalline del campione tal quale (b) , del campione cristallizzato dal fuso non isotermicamente (c) e del campione cristallizzato dal fuso isotermicamente (d). Sono indicati anche i valori degli indici di disaccordo R. Il confronto tra lo spettro di fibra calcolato e quello sperimentale è riportato in Figura 3. Figura 3. Spettro di fibra dell’icisPED14 e suo profilo equatoriale: calcolato (a); sperimentale (b). La proiezione lungo l’asse c del contenuto della cella elementare dell’icisPED14 è riportato in Figura 4.

Figura 4. Proiezione lungo l’asse c del contenuto della cella elementare dell’icisPED14. I gruppi etilici laterali sono riportati in grigio. Conclusioni L’accordo della struttura proposta con i dati sperimentali è molto buono per il campione tal quale, ma non del tutto soddisfacente per il campione ottenuto per cristallizzazione non isoterma dal fuso. Per quest’ultimo campione i due riflessi a 2θ ~ 20° appaiono mal risolti, come se il campione non mostrasse un alto grado di ordine. A riprova, il campione cristallizzato isotermicamente dal fuso a 70°C , cioè a basso sottoraffreddamento, dà luogo ad uno spettro di polveri simile a quello del campione tal quale e perciò in ottimo accordo con quello calcolato. Questo suggerisce che i campioni ottenuti durante cristallizzazioni lente siano più ordinati rispetto a quelli cristallizzati velocemente. Probabilmente, sebbene il gruppo spaziale P212121 preveda la presenza di catene isochirali, durante la cristallizzazione si formano domini sia di cristalli costituiti da catene aventi tutte una chiralità che di cristalli con catene aventi tutte la chiralità opposta. Nelle cristallizzazioni lente tali domini sono più grandi mentre nelle cristallizzazioni veloci i domini sono più piccoli per cui il campione appare più disordinato ai raggi X. La struttura cristallina dell’icisPED14 è molto simile a quella del cis-l,4-poli(1,3-pentadiene) isotattico1,2 [icisPPD14], la cui catena ha gruppi metilici (invece che etilici) come gruppi laterali. Infatti, la struttura dell’icisPPD14 è caratterizzata anch’essa da catene con simmetria s(2/1) e da una cella elementare ortorombica con assi a=9.49Å, b=6.07Å, c=8.17Å, nel gruppo spaziale P212121 . I valori degli assi c delle celle dei due polimeri sono molto simili, mentre il più alto valore di b dell’icisPED14 è dovuto sia al maggiore ingombro del gruppo laterale che all’allineamento delle catene in filari lungo b. I filari, a differenza con l’icisPPD14, sono disposti nella struttura del l’icisPED14 quasi parallelamente e possono bene impacchettarsi per cui per questo polimero l’asse a risulta leggermente più corto rispetto a quello dell’icisPPD14. Riferimenti 1. H. Sun, J. Phys. Chem. B, 102, 7338 (1998). 2. M. Purevsuren, G. Allegra, S. V. Meille, A. Farina, L. Porri,G. Ricci, Polymer J., 5, 431 (1998) 3. R. Napolitano, Macromol Theory Simul., 15, 614 (2006)

R=13%

R=21%

R=14%

5 10 15 20 25 30 35 40

Intensity

2θ (deg)

a

c

d

b

a

b

10 15 20 25 30 35 40

Intensity

2θ (deg)


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EFFECT OF THE SILANIZATION AND ANNEALING ON THE MORPHOLOGY OF THIN P3HT LAYER ON SILICON OXIDE

G. Scavia1, W. Porzio2, S. Destri2 1Istituto di Struttura della Materia-CNR Area di ricerca di Montelibretti Via Salaria, Km. 29,300 -00016 Monterotondo RM, Italy; e-mail [email protected]

2Istituto per lo Studio delle Macromolecole- CNR, via E. Bassini 15, 20133 Milano

Introduction Conjugated polymers such as P3HT are becoming a valid alternative to amorphous silicon for the active layer in thin film transistors for large area and low cost applications [1]. The most significant electronic parameter defining the P3HT/Silicon oxide performance in the OFET device, i.e. electron mobility, have been known to be related to the P3HT crystalline structure of P3HT layer, mainly the first layers close to the interface [2-3]. A method to modulate and consequently optimize the interaction between the polymer/oligomer layer and the silicon oxide substrate for a better electron mobility consists in the functionalization of the substrate with reactive alkyl substituted silane molecules (SiX). The insertion of the resulting SAM layer in between the SiOx and P3HT has been demonstrated to improve the electrical performance of the OFET in terms of parallel electron mobility mainly for the layers immediately close to the interface[2,5]. Even though this silanization has been widely studied from an electronic point of view, little has been done in order to understand the structural modifications onto the first P3HT layers (<5nm) leading to such electronic improvement. For this purpose, in parallel with XRD measurements, AFM measurements onto the thin layers closer to the silanized surface could give a ‘top view’ of the first P3HT layers interacting with the interface. This could be useful for defining a structural model at the interface and for establishing a correlation between structure and electron mobility. The objective of this study is first of all to extrapolate from AFM morphological images of thin P3HT layer, in parallel with XRD measurements, a plausible model of interaction between SAM and P3HT molecules and secondly to see the effect of a change of silanizer SAM property (alkyl length, polarity) onto the corresponding P3HT morphology. Results AFM studies show that a significant change in morphology occurs for silanized substrate. The straight filaments, typical of the P3HT/Silicon system, change into bent filaments more compact and organized for a silicon oxide silanized with a silane molecule SiX with X=Me substrate and then are substituted by needle like structures for SiX=C18(OTS). For a X of intermediate length, a coexistence of the two morphologies occurs. From a detailed analysis two models can be used to explain in molecular terms the two morphologies: for filament the well known π-π interaction model in which P3HT chains are edgewise onto the substrate and for the needle-like case, typical of the long chains, a lamellar organization in which P3HT lies parallel to the substrate. A mechanism of interaction explaining this

‘switch’ from perpendicular filament-like configuration to parallel needle-like configuration is proposed.

Figure1: Comparison between the AFM phase image of P3HT/SiMe2(a) and the corresponding phase image for P3HT/OTS (b). The effect of the post-annealing is presented for all the considered silanizations. The post-annealing at temperatures close to the polymer melting point (190°C) 1results efficient in improving the P3HT organization into ordered structures. Figure 2 shows the post-annealing effect for P3HT deposited on SiMe.

Fig. 2 a) height image of the non-annealed P3HT/SiMe layer (corresponding phase image comparable with the height image) b) phase image of the same layer annealed at 190°C. The higher scale for a) has been used in order to evidence the filament morphology sensibly less dense than in b). Conclusion The effect of the silanization onto the P3HT thin layer morphology deposited by spin coating has been investigated with AFM. Results indicate that by increasing the size of alkyl substituent in the silanizer molecule, a corresponding change in morphology occurs from filament-like to needle-like. The first morphology corresponds to a well known model in which P3HT are disposed perpendicular to the surface thus forming filaments. On the contrary, the needle like morphology could be hypothesised to correspond to a configuration in which the P3HT molecules lie horizontally to the surface thus forming thin lamellar structures.


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Acknowledgements We thank the CARIPLO foundation for partial support by PROTEO project. References [1] C.D. Dimitrakopoulos, P.R.L. Malefant Adv. Mater. 14, 99 (2002)

[2] H.Sirringhaus,N.Tessler,R.H.Friend Science 280, 1741 (1998) [3] R.A.Street, J.E.Northrup, A.Salleo Phys.Rev.B 71, 165202 (2005) [4] R.J. Kline, M.D. Mcgehee, M.F. Toney, Nature Materials 5, 222 (2006) [5] H.Kim, Y.Jang, Y.D.Park K.Cho Macromolecules 39, 5843 (2006)


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TRANSIZIONI DI FASE INDOTTE DA SFORZI TENSILI IN COPOLIMERI SINDIOTATTICI PROPENE-1-BUTENE AVENTI CARATTERISTICHE

ELASTOMERICHE

L. Caliano, C. De Rosa, M. Corradi, F. Auriemma, G. Talarico Università degli Studi di Napoli “Federico II”, Dipartimento di Chimica, Complesso di Monte Sant'Angelo via Cintia, 4 80126 Napoli, Italy; e-mail: [email protected] Introduzione Nella presente comunicazione verranno illustrati i risultati ottenuti in recenti studi condotti nel nostro gruppo di ricerca su polimeri semicristallini prodotti da catalisi metallocenica. Tali studi sono intesi a chiarire le relazioni tra le proprietà fisiche e meccaniche, l’organizzazione strutturale della fase cristallina di tali sistemi e la microstruttura (grado di stereoregolarità, presenza di unità comonomeriche) delle catene. Queste indagini hanno consentito di identificare una nuova classe di elastomeri termoplastici a base di poliolefine caratterizzati da elevata rigidità dovuta alla presenza di una frazione non trascurabile di fase cristallina. Un esempio interessante di tale classe di elastomeri è rappresentato dai copolimeri sindiotattici propene-1-butene (sPPBu). I copolimeri sPPBu mostrano cristallinità in tutto l’intervallo di composizione e sono caratterizzati da temperature di fusione (Tm ) e transizione vetrosa (Tg) decrescenti all’aumentare del contenuto di 1-butene (Tabella 1). Tabella 1: Composizione, temperature di fusione (Tm) e di transizione vetrosa (Tg) e pesi molecolari dei campioni di copolimeri sPPBu analizzati.

a) Catalizzatore: Ph2C(CpFlu)ZrCl2/MAO; Condizioni operative: T=10°C; P=1 atm. La caratterizzazione strutturale di copolimeri sPPBu è riportata in letteratura per campioni non orientati 1, 2. In particolare, è stato dimostrato che l’aumento del contenuto di unità 1-butene comporta la variazione delle

dimensioni degli assi a e b della cella elementare dai valori caratteristici del polipropilene sindiotattico (sPP) a quelli del polibutene sindiotattico (sPB) (Figura 1). Questo risultato costituisce una evidente prova dell’inclusione delle unità 1-butene nella fase cristallina.

Figura 1: Variazione degli assi a (A) e b (B) della cella elementare e della distanza di Bragg (B) all’aumentare del contenuto di butene. Nella presente comunicazione si propone uno studio volto a chiarire la relazione tra le proprietà elastiche di fibre di copolimeri sPPBu in tutto l’intervallo di composizione e le trasformazioni strutturali della fase cristallina indotte dallo stiro. Risultati Gli studi strutturali hanno dimostrato che per bassi contenuti di butene (< 4mol%) l’elasticità di fibre dei copolimeri sPPBu è analoga a quella dell’omopolimero polipropilene sindiotattico (sPP)3 ed è, pertanto, associata ad una transizione di fase reversibile da una forma cristallina con catene in conformazione trans-planare che si trasforma in una forma cristallina più stabile caratterizzata da catene in conformazione elicoidale (Figura 2).

Campioni a 1-butene (mol%) Tm (°C) Tg (°C) Mv×10-5 sPPBu1 3.2 138 0 3.5 sPPBu3 6.7 123 -0.1 3.4 sPPBu4 11.2 110 -4 3.1 sPPBu5 13.6 108 -3 4.1 sPPBu6 18.2 100 -5 2.8 sPPBu7 31.5 85 -7 2.7 sPPBu8 37.9 71 -8 2.0 sPPBu9 51.7 70 -10 1.9 sPPBu10 52.1 64 -11 1.8 sPPBu11 69.9 57 -14 1.4 sPPBu12 89.0 54 -17 2.2

0 10 20 30 40 50 60 70 80 90 10014,014,515,015,516,016,517,0

a (Å

)

mol % 1-butene

A

0 10 20 30 40 50 60 70 80 90 1005,55,65,75,85,96,06,1

5,55,65,75,85,96,06,1

110010

d (Å)

b (Å)

mol % 1-butene

� Asse b � Distanza di Bragg B


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Figura 2: Figure di diffrazione dei raggi X di una fibra del copolimero sPPBu1 con 3.2 mol% di butene stirata al 600% di deformazione e successivamente rilassata. L’aumento del contenuto di 1-butene sfavorisce l’ottenimento per stiro di forme polimorfe con catene in conformazione trans-planare, per cui in cicli consecutivi di stiro e rimozione della tensione non si verifica alcuna trasformazione strutturale della forma cristallina più stabile del sPP con catene in conformazione elicoidale. Nei copolimeri con un contenuto di butene > 70 mol% si ottiene ad elevate deformazioni una nuova forma cristallina del sPB in cui le catene in conformazione elicoidale sono impacchettate in una cella elementare secondo una struttura ortorombica B-centrata che si trasforma nella struttura ortorombica C-centrata più stabile per effetto della rimozione della tensione (Figura 3).

Figura 3: Figure di diffrazione di raggi X di una fibra del copolimero sPPBu11 con 69.9 mol% di butene stirata al 400% di deformazione e successivamente rilassata e corrispondenti modelli di impacchettamento delle catene. La registrazione di cicli ripetuti di stiro e rimozione della tensione ha evidenziato il buon recupero elastico di fibre di campioni ad alto contenuto di butene (Figura 4).

0 10 20 30 40 50 600246810121416

sPPBu11

sPPBu9

ε %

σ (M

Pa)

Figura 4: Curve sforzo-deformazione registrate a Temperatura ambiente per fibre dei copolimeri sPPBu9 e sPPBu11 (51.7 e 69.9 mol% di 1-butene, rispettivamente) durante cicli consecutivi di stiro e rimozione della tensione in accordo con la direzione delle frecce. Conclusioni La caratterizzazione strutturale di fibre di copolimeri sindiotattici propene-1-butene ha evidenziato che per bassi contenuti di butene (< 4mol%) l’elasticità è dovuta ad una transizione reversibile cristallo-cristallo in cui catene in conformazione elicoidale si trasformano in catene in conformazione trans-planare per effetto dello stiro. L’aumento del contenuto di 1-butene impedisce la suddetta transizione stabilizzando le catene in conformazione elicoidale. Per contenuti di 1-butene superiori a 70 mol% è stata identificata una nuova forma cristallina del sPB in cui le catene in conformazione elicoidale sono impacchettate in una struttura ortorombica B-centrata che transisce verso la struttura ortorombica C-centrata più stabile del sPB in seguito alla rimozione della tensione. Il comportamento elastico di fibre di questi copolimeri permane anche ad elevati contenuti di butene per cui si può affermare che per bassi contenuti di butene (< 4mol%) l’elasticità è di natura entalpica in quanto connessa alla transizione cristallo-cristallo analoga a quella osservata per l’ sPP, mentre per alti contenuti di butene (> 60 mol%) è di natura entropica come per gli elastomeri tradizionali. Tuttavia la presenza non trascurabile di una fase cristallina garantisce in tali materiali un grado di rigidità superiore a quello osservato in elastomeri convenzionali. Riferimenti 1. C. De Rosa, G. Talarico, L. Caporaso, F. Auriemma, O. Fusco, M. Galimberti, Macromolecules, 31, 9109 (1998). 2. C. De Rosa, F. Auriemma, I. Orlando, G. Talarico, L. Caporaso, Macromolecules, 34, 1663 (2001). 3. C. De Rosa, F. Auriemma, Prog. Polym. Sci., 31, 145-237 (2006).

ε = 600%

Rilassamento Elica

trans-planare

Stiro

ε =400 %

Rilassamento Stiro

Elica Elica


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RUOLO DEI DIFETTI COSTITUZIONALI NEL COMPORTAMENTO ELASTICO DEL POLIPROPILENE SINDIOTATTICO

M. Corradi, C. De Rosa, F. Auriemma, G. Talarico, G. Giusto, R. Di Girolamo, L. Caliano

Università degli Studi di Napoli “Federico II”, Dipartimento di Chimica, Complesso di Monte Sant'Angelo via Cintia, 4 80126 Napoli, Italy; e-mail: [email protected] Introduzione In questo lavoro è stato analizzato l’effetto dell’introduzione di difetti costituzionali sulle proprietà fisiche del polipropilene sindiotattico (sPP), e in particolare sulle sue proprietà elastiche. Questo peculiare comportamento elastomerico, insolito per un polimero ad elevata cristallinità e con una temperatura di transizione vetrosa relativamente elevata (Tg ~ 0°C), è associato ad una transizione di fase reversibile che fornisce un contributo entalpico al processo di ritorno elastico. L’analisi strutturale ha infatti mostrato che, quando la tensione di stiro viene rimossa, la forma III con catene in conformazione trans-planare, presente in fibre orientate di sPP, si trasforma nella forma II più stabile, caratterizzata da catene in conformazione elicoidale. Si osserva, contemporaneamente, un recupero totale delle dimensioni iniziali delle fibre.1 Obiettivo di questo lavoro è lo studio delle proprietà fisiche e meccaniche di copolimeri sindiotattici del propilene per determinare in che modo la presenza dei difetti costituzionali, rappresentati da unità comonomeriche, influenzi il polimorfismo del sPP e le sue proprietà fisiche ed elastiche. E' infatti atteso che la presenza di comonomeri influenzi le transizioni strutturali tra le forme cristalline con catene in conformazione elicoidale e quelle con catene in conformazione trans-planare, responsabili in parte dell’elasticità del sPP, e che quindi possa determinare una variazione delle proprietà elastiche. Risultati Sono stati preparati campioni di copolimeri sindiotattici del propene con etilene (sPP-C2), 1-butene (sPP-C4), 1-esene (sPP-C6), 1-ottene (sPP-C8), 1-dodecene (sPP-C12), 1-octadecene (sPP-C18) e 4-metil-1-pentene (sPP-4M1P). Tutte le polimerizzazioni sono state condotte a 10 °C utilizzando come catalizzatore Ph2C(Cp)(Flu)ZrCl2 attivato con MAO, che garantisce elevate masse molecolari ed una distribuzione random delle unità comonomeriche. La variazione della microstruttura, associata all’introduzione dei difetti costituzionali, provoca un cambiamento graduale delle proprietà fisiche dell’omopolimero, quali ad esempio la temperatura di fusione (Figura 1) e la temperatura di transizione vetrosa (Figura 2). In particolare, la diminuzione della temperatura di transizione vetrosa è una caratteristica fondamentale per ampliare il campo di applicazioni di questi materiali.

0 10 20 30 40 50 60 70 80 90 100405060708090100110120130140150160 sPP

sPP-C2 sPP-C4 sPP-C6 sPP-C8 sPP-C12 sPP-4M1P

T m (°

C)mol % comonomero

Figura 1: Temperatura di fusione (Tm) dei campioni di copolimero analizzati in funzione della concentrazione di comonomero.

0 10 20 30 40 50 60 70 80 90 100

-25

-20

-15

-10

-5

0

5

10 sPP sPP-C4 sPP-C6 sPP-C8 sPP-C12 sPP-C18 sPP-4M1P

T g (°

C)

mol % comonomero

Figura 2: Temperatura di transizione vetrosa (Tg) dei campioni di copolimero analizzati in funzione della concentrazione di comonomero. Nella Figura 3 sono mostrate le curve sforzo-deformazione di campioni dei copolimeri contenenti piccole quantità di comonomero (2-4 mol%). Questi dati indicano che tutti i campioni presentano interessanti proprietà meccaniche quali buona duttilità e alto modulo. Tutti i campioni inoltre presentano un discreto comportamento elastomerico indipendentemente dalla composizione; fibre orientate degli stessi campioni presentano ottime proprietà elastiche, e valori del modulo di Young maggiori rispetto a quelli dei campioni non orientati.


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0 100 200 300 400 500 600 700 8000

5

10

15

20

250 100 200 300 400 500 600 700 800

0

5

10

15

20

25

300 500 1000 1500 2000 2500

0

5

10

15

20

25

30

4.0% 4M1P

sPP2.7% C12

Strain (%)

Stress

(MPa

)

sPP

3.1% C83.0% C6 3.2% C4

Stress

(MPa

)2.6% C2

sPP

Stress

(MPa

)

Figura 3: Curve sforzo-deformazione di film pressofusi di campioni di copolimeri sindiotattici del propilene. Per ogni campione è riportato il contenuto in mol% di comonomero. Lo studio strutturale ha dimostrato che la presenza di difetti costituzionali influenza notevolmente il polimorfismo del sPP, anche in campioni orientati. Fibre dei copolimeri con basso contenuto di comonomero (≤ 5 mol%) presentano un comportamento molto simile a quello mostrato dall’omopolimero. La forma I elicoidale presente nel campione non stirato si trasforma nella forma III, con catene in conformazione trans-planare, già per bassi rapporti di stiro. In seguito al rilascio della tensione la forma III si trasforma nella forma II elicoidale, ed in parte nella forma disordinata mesomorfa con catene ancora in conformazione trans-planare Nel caso dei copolimeri con l’etilene, la presenza del comonomero provoca la stabilizzazione delle forme cristalline con catene in conformazione trans-planare;2 in tutti gli altri copolimeri è stato invece osservato che l’aumento della concentrazione di comonomero provoca

una stabilizzazione delle forme cristalline con catene in conformazione elicoidale. Per elevate concentrazioni di comonomero quindi, non si osserva alcuna transizione polimorfa dopo rimozione della tensione. Ne consegue che l’entità della trasformazione polimorfa che si osserva rimuovendo la tensione di stiro, diminuisce progressivamente all’aumentare della concentrazione di comonomero, a causa della stabilizzazione delle forme cristalline ottenute durante lo stiro. Pertanto solo per i campioni con più basso contenuto di comonomero il ritorno elastico può essere associato ad una trasformazione polimorfa come accade nell’omopolimero. Poiché tutti i campioni presentano comportamento elastico, l’elasticità presenta diversi contributi in funzione della composizione e della cristallinità. Conclusioni Questo studio ha mostrato che il comportamento elastomerico dei copolimeri sindiotattici del propilene è il risultato di due diversi contributi. L’introduzione di unità comonomeriche in catena provoca infatti una variazione delle proprietà fisiche dell’omopolimero, e influenza il suo complesso polimorfismo. Per campioni con basso contenuto di comonomero ed elevata cristallinità, l’elasticità ha principalmente natura entalpica (come nell’omopolimero sPP), dovuta alla metastabilità della forma III che si trasforma nelle forme elicoidali più stabili durante il rilassamento. Il contributo entalpico svolge un ruolo sempre meno importante all’aumentare della concentrazione di comonomero e, per i campioni con alti contenuti di unità comonomeriche, l’elasticità è puramente di natura entropica ossia dovuta alle transizioni conformazionali delle catene della fase amorfa, tra le conformazioni estese, che si formano nello stiro, e le conformazioni random coil (entropicamente favorite) che si riformano dopo rimozione della tensione. Per campioni con composizioni intermedie l’elasticità è dovuta ad entrambi i contributi entalpici ed entropici. Riferimenti 1. C. De Rosa, F. Auriemma, Prog. Polym. Sci., 31, 145 (2006) 2. C. De Rosa, F. Auriemma, Advanced Materials, 17, 1503 (2005)


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MACROCYCLIC ENSEMBLE IN ANIONIC POLYAMIDE 6

L. Ricco1, E. Casazza1, S. Russo1, E. Scamporrino2, P. Mineo2 1Dipartimento di Chimica e Chimica Industriale, Università di Genova, e centro INSTM-NIPLAB, Via Dodecaneso 31, 16146 Genova, Italy; e-mail:[email protected] 2Dipartimento di Scienze Chimiche, Università di Catania, e centro di riferimento INSTM, Viale A.Doria 6, 95125 Catania, Italy

Introduction Recent SEC chromatography data on anionic poly(ε-caprolactam) (PCL) have shown the unexpected presence of a secondary peak at elution times intermediate between the high polymer trace and those of the extractable oligomers. Such peak, coded GR peak, has always been found present in quasi-adiabatic as well as isothermal anionic CL polymerizations, although it has never been evidenced by previous literature. A careful fractionation of the various anionic polymerization products allowed to isolate the GR peak and study this latter by several techniques, in order to identify the nature of its constituents, the parameters that affect their formation and, possibly, the mechanism of their formation. The fractionation procedure has been specifically applied to PCL samples synthesized in isothermal conditions, in order to avoid side products and PCL structural irregularities, typical of CL anionic polymerization when performed under uncontrolled conditions (e.g. quasi-adiabatic polymerization). Indeed, side products can make very difficult the separation of GR peak species from the other products of polymerization as well as the interpretation of experimental results. Results The nature of the fraction corresponding to GR peak has been investigated by SEC and HPLC chromatography, FTIR and NMR spectroscopy, DSC and MALDI-TOF analysis. In particular, SEC chromatography, performed under specific conditions, gave a multimodal chromatogram (Figure 1) corresponding to a mixture of different molecules with peaks characterized by several molar masses. Figure 1: SEC analysis of GR peak of anionic PCL Corresponding results were obtained by DSC analysis that showed the presence of species melting at different, albeit very near temperatures, starting from 170°C and ending at about 210°C. FTIR and NMR spectroscopy suggested that the low molar mass species forming the

GR peak are molecules composed of PCL repetition units. This result has been confirmed by MALDI-TOF analysis, that revealed that the peak is entirely composed of macrorings made of multiples of ε-caprolactam repetition unit, typically from M8 to M20. Moreover, a quantitative interpretation of MALDI spectrum, together with HPLC analysis, gave important information on the distribution of cyclic species from M1 to M11. The distribution was found very different from the expected classical distribution suggested by the Jacobson-Stockmayer theory1, valid only under equilibrium conditions. Namely, when oligomers from CL to pentamer are considered, a regularly decreasing trend as a function of molar mass is observed followed, from M6 on, by an increase up to a maximum corresponding to M9 - M10. Therefore, annealing tests have been performed on our anionic PCL in order to investigate if cyclic species (M1-M20), formed under kinetic control, change their distribution when the thermodynamic equilibrium is reached. Surprisingly, and in contradiction to the Jacobson-Stockmayer theory, MALDI-TOF analysis showed that the molar content of cycles from M8 to M20 maintains the same trend even after annealing at 250°C. In order to explain the formation and the relative concentration of these macrocycles, CL has been anionically polymerized by changing a few parameters of the synthesis. The products of these reactions have been analyzed by SEC chromatography and MALDI-TOF technique, and have shown that the content of the macrocyclic fraction increases as the high polymer fraction (i.e. the polymerization time) increases. On the light of the results obtained so far, a mechanism that could explain the formation of the investigated species has been proposed. Conclusions The previously undiscovered peak, situated between the PLC trace and the cyclic oligomer fraction, is made of macrocycles only, ranging about from the octamer up to M20. These species have been found to grow with the high polymer, and their overall distribution, even in equilibrium conditions, seems to not obey to the Jacobson-Stockmayer theory. A confirmation of the latter results by SEC analysis under suitable conditions is planned. References 1. H. Jacobson, W. H. Stockmayer, J. Chem. Phys., 18, 1600-1606 (1950)


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STUDIO DELL'EFFETTO DELLA VISCOELASTICITA' NEL METODO DEL LAVORO ESSENZIALE DI FRATTURA

C. Marano, R.Rizzieri, M. Rink

Politecnico di Milano, Dipartimento di Chimica, Materiali e Ingegneria chimica “Giulio Natta”, Piazza L. da Vinci, 32, 20133, Milano, Italy; e-mail: [email protected] Introduzione Il metodo del lavoro essenziale di frattura (EWF) [ 1] può essere impiegato per valutare la tenacità di film polimerici ed è basato sulla scomposizione del lavoro di frattura, Wf, in due componenti We e Wp, rappresentative, rispetivamente, dell’energia necessaria alla creazione di nuove superfici e dell’energia necessaria alla deformazione del materiale, con dissipazione di energia, nell’intorno della zona di frattura. Il principio su cui si basa la tecnica è la misura dell’energia di frattura, Wf, da curve carico-spostamento relative ad una serie di campioni intagliati con differente valore della lunghezza della sezione resistente, L. L’energia specifica di frattura, wf, risulta così essere funzione della lunghezza della sezione resistente L:

w f (=W f

LB ) = we + βw pL eq. 1 Per estrapolazione a L = 0 della retta che interpola i dati sperimentali, si determina il lavoro essenziale we, e dalla pendenza il lavoro specifico non-essenziale di frattura βwp. Il primo termine è l’energia per unità di superficie dissipata nei processi che portano alla formazione di nuove superfici, mentre βwp è l’energia per unità di volume dissipata nell’intorno della zona di frattura. β è un fattore di forma che dipende dalla geometria del volume di materiale coinvolto. E’ stato osservato [2] che se nella valutazione di Wf, lo spostamento è misurato ai margini della zona di processo invece che ai margini dell’intero campione, come suggerito dal protocollo [1], il termine βwp diminuisce, mentre we non varia. Questo risultato corrisponde ad una diminuzione, dipendente da L, dell’energia specifica di frattura e suggerisce che la deformazione del materiale al di fuori della zona di processo dissipi energia che contribuisce al termine wf attuale. In questo lavoro l’energia di deformazione dissipata dal materiale al di fuori della zona di processo è stata misurata sperimentalmente allo scopo di confrontarla con le variazioni osservate in wf misurato con il metodo EWF. Parte sperimentale Materiali Il materiale studiato in questo lavoro è un film estruso di un copolimero a blocchi etilene-propilene (nome commerciale Moplen EP 310D) fornito da BASF (D), dello spessore medio B pari a 100 µm, caratterizzato da una temperatura di fusione di 164°C ed un’entalpia specifica di fusione di 71 J/g (proprietà misurate in un calorimetro Mettler TA alla velocità di riscaldamento di 10°C/min). Le prove sono state condotte utilizzando provini intagliati (DEN) di larghezza W = 35 e altezza H = 35 mm (Fig. 1), ricavati dal film, mediante un bisturi, in modo che la direzione di estrusione del film coincida

con la direzione di applicazione del carico. Gli intagli sono stati realizzati per penetrazione di una lama (raggio di curvatura all’apice pari a circa 13 µm) e tali da avere lunghezza della sezione resistente L = (W – a) pari a circa 5, 7, 13 e 15 mm (a = lunghezza dell’intaglio) (Fig. 1). Dettagli sperimentali Prove di trazione sono state eseguite, a temperatura ambiente, mediante un dinamometro INSTRON 1121, alla velocità di spostamento v = 10 mm/min, Su ciascun provino sono stati disposti 10 “markers” in maniera simmetrica rispetto alla sezione resistente, come illustrato in figura 1, a partire da una distanza da questa pari a L/2, nell’ipotesi plausibile [1] che la zona di processo abbia un’estensione pari ad L. La distanza, d, fra due markers successivi è pari ad 1 mm. Mediante un videoestensometro sono state osservate nel tempo le coordinate (xi, yi) dei singoli markers. È stato possibile misurare, in questo modo, lo spostamento, ∆centrale, ai margini della zona di processo, di estensione iniziale pari ad L, come suggerito in [2] e definito come

∆centrale = [(y5 – y6) – L] eq.2. Inoltre, dalle variazioni delle distanze, ∆i, tra coppie di markers adiacenti

∆i = [(yn – yn+1) – d] eq.3 è stato possibile misurare le deformazioni locali εi = ∆i/d al di fuori della zona di processo.

Figura 1: schema e dimensioni del provino utilizzato nel test EWF Risultati Determinazione del lavoro essenziale di frattura In Figura 2 il lavoro specifico di frattura, wf, (wf = Wf/BL) è diagrammato in funzione della lunghezza della sezione resistente, L. Nel calcolo dell’energia di frattura Wf, lo spostamento, ∆, è stato misurato: (a) al contorno (∆TOT = v · t); (b) ai margini della zona di processo, ∆centrale (eq.2). In tabella 1 sono riportati i valori del lavoro essenziale di frattura, we, e del lavoro non essenziale di frattura, βwp, nei due casi a e b. Come già osservato in [2], solo il lavoro non essenziale di frattura, βwp, dipende dal metodo di misura dello spostamento ∆. Questo risultato si spiega ammettendo che, nella misura del lavoro di frattura wf attraverso lo spostamento al contorno (metodo


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a) si include un contributo associato alla deformazione del provino al di fuori della zona di processo, non considerato nel metodo b. Tale contributo (∆wf = wf_a - wf_b) aumenta all’aumentare della lunghezza della sezione resistente, L, ed è attribuibile ad una deformazione viscoelastica del materiale. Infatti, in questo lavoro è stato verificato che non vi è deformazione permanente, cioè plastica, nella deformazione residua misurata nel provino dopo frattura, al di fuori della zona di processo. Nell’ipotesi di una deformazione viscoelastica nel volume di materiale V al di fuori della zona di processo dovrebbe, pertanto, risultare:

∆wf · BL = wve· V = (∫σ dε) · V eq. 4 dove σ = P/(WB) è lo sforzo applicato al di fuori della zona di processo, in cui si misura la deformazione ε.

0

50

100

150

200

250

0 5 10 15

metodo ametodo b

Lavo

ro spe

cifico

di fra

ttura,

wf (k

J/m2 )

Lunghezza sezione resistente, L (mm)

we

βwp

Figura 2: lavoro specifico di frattura, wf, in funzione della lunghezza, L, della sezione resistente: a_lo spostamento è stato misurato al contorno; b_lo spostamento è stato misurato ai margini della zona di processo. Tabella I spostamento we (kJ/m2) βwp (MJ/m3) (a) ∆TOT 38.3 11.5 (b) ∆centrale 36.9 10.8

Determinazione dell’energia di deformazione al di fuori della zona di processo In Figura 3 è riportato, in funzione del tempo, l’andamento della deformazione locale εi = ∆i/d misurata al di fuori della zona di processo durante una prova di frattura. I dati riportati a titolo di esempio sono relativi ai campioni con lunghezza della sezione resistente, L, pari a 7 mm. Come prevedibile, in ragione della geometria del campione, la deformazione misurata raggiunge valori significativamente più elevati in prossimità della zona di processo (ε4_5, ε6_7) rispetto alle estremità del provino (ε1_2, ε9_10). In figura 4 sono riportate curve isocrone delle deformazione locali, εi, in funzione della distanza dalla sezione resistente, h. Si può osservare che il gradiente di deformazione che si genera al di fuori della zona di processo (h ≥ 3.5 mm nel caso considerato), è relativamente piccolo nei primi istanti della prova e successivamente per effetto di un incremento della deformazione più significativo in prossimità della zona di processo.

0

1

2

3

4

5

6

0 5 10 15 20 25 30 35 40

Defor

mazio

ne, ε

(%)

tempo (s)

ε4-5

, ε6-7

ε3-4

, ε7-8

ε1-2

, ε9-10

Figura 3: deformazione locale, in funzione del tempo (L=7 mm) Dai risultati ottenuti sono state valutate le energie dissipate nel volume del campione al di fuori della zona di processo. A questo scopo tale volume è stato suddiviso in parti di volume Vi, in cui determinare l’energia dissipata localmente (∫σ dε)ι tenendo conto dello sforzo locale applicato, σi = F/Si. E’ stato trovato un buon accordo tra l’energia specifica dissipata così valutata e la differenza ∆wf = wf_a - wf_b tra i valori dell’energia specifica di frattura determinata secondo i due metodi a e b precedentemente descritti.

h

0

1

2

3

4

5

6

0 0.2 0.4 0.6 0.8 1

2.3 s 6.1 s 13.6 s 35.0 s

Defor

mazio

ne, ε

, (%)

Distanza normalizzata dalla sezione resistente,

0

10

20

0 10 20 30 40

F (N)

t (s)

Figura 4: curve isocrone della deformazione locale (H=35 mm, L=7 mm). Conclusioni In questo studio è stato verificato che nella prova del lavoro essenziale di frattura, per il materiale considerato, vi è dissipazione di energia al di fuori della zona di processo, per effetto di una deformazione viscoelastica del materiale. Come già riportato in altri lavori, è stato osservato che il termine del lavoro essenziale di frattura, we, non risente significativamente del metodo di misura del lavoro specifico adottato che invece influisce sul valore del lavoro specifico non essenziale, βwp. Riferimenti 1. E. Clutton, ESIS Test protocol Essential Work of Fracture (revised by D R Moore 2002). 2. M.LI. Maspoch, G. Rodriguez, J. Gamez-Perez, A. Gordillo, O.O. Santana, 4th ESIS TC4 CONFERENCE 11-14 September 2005, Les Diablerets (CH)


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SODIUM ALGINATES AND GLYCEROL BASED FILMS: INFLUENCE OF DIFFERENT POLYMERIC MICROSTRUCTURES ON CHEMICAL-PHYSICAL

PROPERTIES OF THE FILMS

M. Malinconico1, B. Immirzi1, E. Di Pace1, M. Avella1, G. Santagata1, G. Impallomeni2 1 ICTP - CNR, Via Campi Flegrei, 34 – 80078 Pozzuoli, Napoli, Italy; e-mail: [email protected]

2 ICTP - CNR, Viale A. Doria, 6 – 95125 Catania, Italy Introduction Nowadays, in agricultural applications such as mulching and solarization techniques, synthetic polymers are commonly used; the films obtained by them present excellent mechanical features and restrained prices, but considering the extended ecological and environmental concern, they both don’t satisfy the increasing demand of biodegradability, and they contribute to the problems related to their disposal like special litter. To minimize this problem and to reduce the strong dependence on petroleum, we can identify biodegradable polymers as new starting materials to be used in film preparation. To this purpose, natural polysaccharides are very interesting resources for film applications, because of their wide availability and renewability (1, 2,). The aim of this study is to use sodium alginate as a base (matrix) of film preparation. Sodium alginate is a polysaccharide coming from brown seaweeds, with the peculiar ability to undergo irreversible cross-linking in the presence of divalent cations such as calcium, becoming insoluble in water and, hence, resulting suitable for uses on wet soils, Fig 1

Fig 1 Egg-box structure Unfortunately, films of pure alginate do not show good mechanical properties in a useful range; in fact they are characterized by a very high modulus and a very low elongation at rupture. The addition of glycerol like plasticizer is necessary to reduce the films fragility, particularly upon their cross-linking. In this paper we studied the chemical-physical properties of biodegradable films obtained by water blending and casting sodium alginates with increasing amounts of glycerol plasticizer. Two alginates with different molecular weights and different M/G ratios have been used and mechanical, thermal and morphological properties of the films performed have been studied. Two types of alginate were used: one of high guluronic acid content (Protonal LF 20 alginate with the identification code “Ap” and one of low guluronic acid content (Scogin LV alginate with the identification code “Ar”), Glycerol was provided by Fluka and used as received. The compositions of the blends with their identification codes are reported in Table 1.

Table 1. Formulations Code

wt % Protonal LF 20 alginate

wt % Scogin LV alginate

wt % glycerol

Ap 100 - 0 ApG1 67 - 33 ApG2 55 - 45 ApG3 50 - 50 Ar - 100 0 ArG1 - 67 33 ArG2 - 55 45 ArG3 - 50 50 On the obtained films, thermal analysis (DSC, DMTA) mechanical properties and morphological analysis (SEM) have been performed Results and discussion The different thermal performance of the two neat polymers and of the blends suggests the hypothesis of a strict correlation between the observed properties and the polymeric microstructures. Comparing DSC curves of Ap and Ar systems Figg 2, 3, it finds out that the endothermic peaks related to maxima water losses are located at different range of temperatures; in fact for Ar system the temperature range is 80° C-100° C, while the interval for Ap system is 110° C-130° C. Comparing DMTA curves (not reported) of Ap and Ar based systems, it is evident the diverse plasticizing action of glycerol on polymeric matrixes; in ApG2 and ApG3 blends, the excess of plasticizer provides a drastic drop of modulus and a decreasing of polymer Tg, while in Ar system only the ArG3 blend shows an early plasticizing effect. The different thermal behaviours could be explained on the base of the micro structural features of Ap and Ar polymers. In Ap polymer the higher value of Mw is indicative of a longer average chain length, and the higher fraction of guluronic acid is accountable of a buckled and folded structure in which a plasticizer is probably entrapped quite firmly; in fact, the addition of 33% of glycerol causes a clear plasticizing effect of the polymer, as shown by the DSC and the log E’ curves Because of this stronger interaction, the water can release the polymeric binding sites at lower temperatures with respect to the neat polymer. On the other hand, DSC curves of ApG2 and ApG3 blends evidence a rising of water evolution temperature this occurrence seems to indicate that at 33 % glycerol content the accessible interactive sites of the polymer have been already


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saturated and increased amounts only goes to interact with water. In Ar polymer, the lower value of Mw is indicative of shorter average chain length and the higher fraction of mannuronic acid accounts of a greater overall flexibility. In this network the water is not strongly retained; in fact, evaluating the ranges of temperature in which the endothermic peaks of Ap and Ar systems occur, it is possible to note a shift towards lower values of water evolution temperature for Ar system.

Fig 2 DSC of ApG system

Fig 3 DSC of ArG system In Tab 2 mechanical properties of ApG system are reported Table 2. Mechanical parameters

Sample Young Modulus (MPa) ± 10% σb (MPa) ± 20% εb (%) ± 20% Ap 2326 44.9 2.13 ApG1 431 20.1 9.40 ApG2 74 6.0 17.7 ApG3 17 5.3 32.2 Ar 3456 82.7 3.75 ArG1 1135 21.93 5.05 ArG2 1054 21.11 5.05 ArG3 626.4 18.0 8.70 Both the systems show a decreasing of Young Modulus going from the neat polymer to the blends In Ap system, moving from the neat polymer to the blends, it is

possible to note a decrease of Young Modulus since the plasticizing effect of glycerol occurred. As concerning the blends, moving from ApG1 to the other blends, a notable drop in the elastic modulus and stress at break are detected. Such a result may be attributed to the excessive amounts of glycerol introduced in the blends and to the resulting change of chemical interactions among polymer, glycerol and water as discussed before. In Ar system, passing by the neat polymer to ArG1 and ArG2 blends, modulus and stress at break slowly decreases and, only in ArG3 composition, a considerable plasticizing effect is observable, in accordance with the hypothesis previously suggested. Conclusions From the above results the following conclusions may be obtained. The plasticizing effect of glycerol on Ap alginates is effectively regulated by mass ratios between the polymers and the glycerol. The properties of blends are strongly related to the molecular structure of the polysaccharides used; in particular high guluronate residues content and high molecular weight represent the correct compromise of micro structural properties able to allow to the glycerol molecules to be entangled in specific binding sites of the polymeric network; As concerning the potential applications of the films obtained, our research is targeted to the preparation of Ap based formulations to be tested on the soil as mulching coatings. For this particular purpose the choice of Ap alginate seems to be more reasonable in view of the mechanical behaviour and of the possibility to tune them with glycerol. In fact, although the recorded mechanical properties of the obtained films are lower if faced with those ones reported for starch based films, nevertheless we must consider that the proposed application of our formulations is substantially different from the application of traditional pre-formed films; in fact the spray methodology doesn’t produce a self-standing film, but a coating supported by the underneath soil. To this aim, work is now in progress. The future target of our study consists in spraying the water solution of polymeric blends on the soil, in this way producing a protective thin membrane (a sort of varnishes) that covers the soil. The spray methodology could represent an absolute innovative practice to prepare films for agriculture. References 1) Immirzi B.; Malinconico M.; Romano G.; Russo R.; Santagata G. J. Mat. Sci. Letters, 22, 1389, 2003. 2) Schettini E., Vox G., Malinconico M., Immirzi B., Santagata G., Acta HorticultureISHS 691, 2005


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WATER DIFFUSION IN POLYIMIDES AS INVESTIGATED BY TIME-RESOLVED FTIR SPECTROSCOPY

P.Musto1, G.Ragosta1, G.Scarinzi1, G.Mensitieri2, M.Lavorgna3

1 Institute of Chemistry and Technology of Polymers, National Research Council of Italy, via Campi Flegrei 34, 80078 Pozzuoli, Italy 2 University “Federico II” of Napoli, Department of Materials and Production Engineering, P.le Tecchio 80, 80125 Napoli, Italy

3 Institute of Composite and Biomedical Materials, National Research Council of Italy, P.le Fermi 1, 80055 Portici, Napoli, Italy

Introduction Polyimides (PI’s) represent an important class of high-performance polymers that exhibit very high Tg, ductility, thermal oxidative stability, and solvent resistance, coupled with relatively low permittivity and dielectric losses up to very high temperatures. For these reasons PI’s find extensive use in microelectronics, photonics, optics and aerospace industry, as well as in separation membrane technology [1]. The PI’s exhibit a complex behaviour in terms of sorption of low molecular weight penetrants. In particular, the diffusion of water has a strong impact on the end-use characteristics of these materials especially in microelectronic applications where it may cause increase in the dielectric constant, metal corrosion of devices, loss of adhesion and mechanical failure. It has been demonstrated that the introduction of fluorinated monomers in the main chain may be used to control the amount of sorbed water and to reduce molecular interactions. However, despite the technological relevance of the problem, the details of the water diffusion mechanism at molecular level are still poorly understood. This mechanism has been investigated in the present study for the case of a commercial polyimide (PMDA-ODA) and two fluorinated polyimides differing for their molecular structure. Several approaches have been used to analyze the time-resolved FTIR data, namely difference spectroscopy (DF), least square curve fitting (LSCF) [2] and two-dimensional correlation spectroscopy (2D-FTIR) [3]. Results Polyimides films were prepared by casting a 10%wt/wt solution in dimethylpyrrolidone (DMP) on a glass plate. Afterwards, the system was subjected to a high temperature curing cycle (up to 300°C) to perform the imidization reactions. The chemical structure of the three investigated polyimides is reported below:

N N

O

O

O

O

On

PMDA-ODA

O

OO

O

ON

C

N

CF3CF3

n

6FDA-ODA O

O O

C

CF3

CF3

O

N

C

N

CF3 CF3

n

6FDA-6FpDA Water vapor sorption and desorption experiments were carried out by using a vacuum tight FTIR cell which allowed to record the transmission spectra of the films exposed to water vapor at constant pressure. Full details of the experimental set-up are reported elsewhere [2]. The experiments were performed at 30°C and at relative pressures p/p0 ranging from 0.1 to 0.80 (where p is the experimental pressure and p0 is the saturation pressure of water at 30°C). Difference spectroscopy has been successfully employed to isolate the spectrum of the penetrant, thus allowing a quantitative determination of its concentration and of the diffusion coefficients. The interacting sites onto the polymer backbone have been identified as the imide carbonyls (see Figure 1).

Figure 1. Difference spectroscopy for a PMDA-ODA film equilibrated at various water vapor activities (spectra labeled B,C,D,E). Spectrum A is related to the dry film; Trace F represents the difference spectrum computed after a sorption/desorption cycle.

B C

D E

A

1800 1600 1400 1200 1000 800 600

--0.2

--0.1

0

0.1

0.2

0.3

1777

1726 1501

1378 1244

1116 1092 882

822 725 604

F

Wavenumber (cm-1)

Abso

rban

ce


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The population of such groups involved in H-bonding with water molecules depends on the experimental conditions and in particular, on the water vapor activity aw. A method has been developed to quantify such population, based on the analysis of the carbonyls stretching vibrations of the imide groups. It has been found that only a fraction (35 %) of the total carbonyl groups is available to interact with water molecules. With respect to the state of aggregation of the penetrant in the polymer matrix, the spectroscopic analysis allowed us to identify two distinct species of H2O molecules. One corresponds to a single H2O molecule interacting with the carbonyl group according to a 1:1 stoichiometry; the other is a self associated aggregate of several water molecules, preferentially dimers. 2D-FTIR spectroscopy effectively enhanced the spectral resolution thus allowing us to identify components undetectable in one-dimensional, frequencies domain. Furthermore, this approach permitted to investigate the dynamic behavior of the system in terms of the relative rate of diffusion of the different water species. It was found that, on sorption, the H2O molecules bound to carbonyls enter the system at faster rate than the self interacting aggregates, while the opposite occurs in the desoption experiment. The diffusion mechanism was rationalized using a model which accounts for an instantaneous, non-linear BET equilibrium between the two water species. Preliminary results on two fluorinated polyimides confirm that, in the whole range of investigated activities, the amount of sorbed water at equilibrium decreases strongly when fluorinated monomers are introduced in the polymer backbone (see Figure 2).

Figure 2. Water vapor sorption isotherms measured by gravimetric methods The strongest reduction of water sorption was achieved when both di-anydride and di-ammine monomers are fluorinated. 2D-FTIR spectroscopy for the 6FDA-6FpDA polyimide evidenced the presence of a component at 3710 cm-1 not identified in the 2D spectra of the other two polyimides (compare Figure 3A and Figure3B). This additional

component has been tentatively assigned to the presence of monomeric water molecules not involved in any H-bonding interaction.

Wav

enum

ber (

cm-1)

Wavenumber (cm-1) 3200 3300 3400 3500 3600 3700 3800 3900

3200 3300 3400 3500 3600 3700 3800 3900 -0.01 -0.008 -0.006 -0.004 -0.002 0 0.002 0.004 0.006 0.008 0.01

Wav

enum

ber (

cm-1)

Wavenumber (cm-1) 3300 3400 3500 3600 3700 3800

3300 3350 3400 3450 3500 3550 3600 3650 3700 3750 3800 -2

-1.5 -1 -0.5 0 0.5 1 1.5 x 10 -3

Figure 3. Asynchronous correlation spectrum of A) PMDA-ODA and B) 6FDA-6FpDA for sorption experiment at p/p0=0.6

Conclusion 2D-FTIR spectroscopy represents a powerful technique to investigate the molecular mechanism of water diffusion in PI’s. Several water species have been identified related to the different polymeric environment. Along with H2O molecules interacting with carbonyl groups and self-associated in aggregates, which are always present in the different PI’s, the high content of fluorine moieties in the 6FDA-6FpDA polymer induces the presence of an additional water specie identified as monomeric water not involved in H-bonding. The various species display different and peculiar dynamic behavior in the sorption and desorption experiments. References 1. Ghosh MK, Mittal KL, editors. Polyimides; Fundamentals and Applications, Marcel Dekker, New York, 1996 2. P.Musto, L.Mascia, G.Mensitieri, G.Ragosta, Polymer, 46,4492 (2005) 3. I.Noda, Appl Spettrosc, 54, 994 (200)

0.0 0.2 0.4 0.6 0.8 1.00

1

2

3

4

PMDA-ODA 6FDA-ODA 6FDA-6FpDA

(Mas

s wate

r/Mas

s dry

polym

er)*1

00

Water vapor pressure, p/p0

A

B


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0 1 0 0 2 0 0 3 0 0 4 0 0 5 0 0 6 0 00

5

1 0

1 5

2 0

2 5

s P P H e 4

s P P H e 3

s P P H e 2

s P P H e 1

Stress

(MPa

)

0 2 0 0 4 0 0 6 0 0 8 0 0 1 0 0 00

5

1 0

1 5

2 0

s P P H e 8s P P H e 5

s P P H e 6

s P P H e 7

Stress

(MPa

)

S t r a i n ( % )

RELAZIONI PROPRIETÁ - STRUTTURA DI COPOLIMERI SINDIOTATTICI PROPILENE-ESENE

G. Giusto, C. De Rosa, F. Auriemma, M. Corradi Università degli Studi di Napoli “Federico II”, Dipartimento di Chimica, Complesso universitario di Monte Sant’Angelo, via Cinthia, 80126, Napoli, Italy; e-mail: [email protected] Introduzione In letteratura è stata già riportata la sintesi ed una caratterizzazione strutturale di copolimeri sindiotattici polipropilene-esene (sPPHe).1,2 Questi copolimeri presentano una discreta cristallinità con temperatura di fusione variabile da circa 150°C (simile a quella del sPP) fino a valori di circa 50°C. In questo lavoro viene riportato uno studio dell’influenza dell’esene sulle proprietà fisiche del sPP, e in particolare sulle proprietà elastiche. L’elasticità del sPP è associata ad una transizione di fase cristallo-cristallo reversibile. La forma III, caratterizzata da catene in conformazione trans-planare, presente in fibre orientate di sPP, si trasforma nella forma II più stabile, con catene in conformazione elicoidale, quando la tensione di stiro viene rimossa. Contemporaneamente si osserva un recupero totale delle dimensioni iniziali delle fibre. In passato è stato suggerito che questa transizione fornisce un contributo entalpico al ritorno elastico del sPP3. Obiettivo di questo lavoro è lo studio delle proprietà fisiche e meccaniche di una serie di campioni di copolimeri sindiotattici propilene-esene per valutare mediante lo studio delle proprietà meccaniche e una caratterizzazione strutturale di fibre orientate per stiro uniassiale, in che misura la presenza di esene influenzi il polimorfismo del sPP e le sue proprietà elastomeriche. In particolare è nostra intenzione determinare se, e in che modo, la presenza dei difetti costituzionali rappresentati dalle unità comonomeriche influenzi le transizioni strutturali tra le forme cristalline con catene in conformazione elicoidale e quelle con catene in conformazione trans-planare, transizioni che nel caso del sPP sono alla base del particolare comportamento elastomerico osservato. Risultati Sono stati sintetizzati campioni di copolimeri sPPHe utilizzando il catalizzatore metallocenico a simmetria Cs Ph2C(CpFlu)ZrCl2, e attivato con MAO. Tutte le polimerizzazioni sono state condotte in toluene a 10 °C. In tabella 1 sono riportate le caratteristiche dei campioni analizzati. Le curve stress-strain (figura 1) hanno mostrato che tutti i campioni presentano interessanti proprietà meccaniche quali buona duttilità, alto modulo e discreto comportamento elastomerico. Fibre orientate degli stessi campioni presentano ottime proprietà elastiche indipendentemente dalla composizione, e valori del modulo di Young maggiori rispetto a quelli dei campioni non orientati.

Tabella 1: Composizione, temperature di fusione Tm e pesi molecolari dei campioni di sPPHe analizzati.

Campioni esene (mol%) Tm (°C) Mv

sPPHe1 1.7 135.8 4.1 ⋅105

sPPHe2 3.0 125.2 5.1 ⋅105 sPPHe3 3.9 111.3 3.1 ⋅105 sPPHe4 5.6 99.8 4.1 ⋅105 sPPHe5 6.4 83.2 3.6 ⋅105 sPPHe6 9.0 68.3 3.2 ⋅105 sPPHe7 11.2 59.0 3.1 ⋅105 sPPHe8 18.8 53.6 3.5 ⋅105

Lo studio strutturale ha dimostrato che l’esene influenza notevolmente il polimorfismo del sPP, stabilizzando le forme cristalline con catene in conformazione elicoidale. Fibre orientate dei copolimeri con basso contenuto di esene (< 2-3 mol%) quando sottoposte a stiro subiscono, già a basse deformazioni, una transizione dalla forma I con catene in conformazione elicoidale alla forma trans-planare mesomorfa pura; per ulteriore stiro a più alti gradi di deformazione la forma mesomorfa si trasforma nella forma III pura con catene ordinate in conformazione trans-planare. In seguito al rilascio della tensione la forma III trans-planare si trasforma nella forma II elicoidale, similmente a quanto accade in fibre di sPP. Figura 1: Curve sforzo deformazione di film pressofusi non orientati di campioni dei copolimeri sPPHe stirati a temperatura ambiente fino a rottura.


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All’aumentare del contenuto di esene il comonomero destabilizza progressivamente la forma III trans-planare, ostacolandone completamente la formazione già in fibre di campioni con concentrazione di esene maggiore del 3 mol%. Al di sopra di tale concentrazione, infatti, la forma III trans-planare non si osserva più e le fibre stirate ad alte deformazioni sono essenzialmente nella forma trans-planare mesomorfa, in miscela con un ammontare di cristalli in forma I elicoidale che aumenta all’aumentare della concentrazione di esene. Quindi la presenza di esene stabilizza sempre di più la forma elicoidale. Per campioni con un contenuto di esene maggiori del 6 mol%, la rimozione della tensione da fibre stirate a varie deformazioni provoca una transizione della forma mesomorfa trans-planare, ottenuta nello stiro, nella forma I elicoidale, indipendentemente dal grado di deformazione raggiunto nello stiro. In tali campioni, contrariamente a quanto accadeva nei copolimeri con basso contenuto di esene, non si è mai osservata la formazione della forma II isochirale in seguito alla rimozione dello stiro. Questo può essere dovuto al fatto che in tutti questi campioni sono sempre presenti nelle fibre stirate cristalli di forma I, in miscela con la forma trans-planare III o mesomorfa, che probabilmente agiscono da nuclei favorendo la trasformazione nella forma I piuttosto che nella forma II. L’entità della trasformazione polimorfa che si osserva rimovendo la tensione di stiro diminuisce progressivamente all’aumentare della concentrazione di esene, dal momento che una sempre minore quantità di forma mesomorfa si ottiene nelle fibre stirate, che sono infatti prevalentemente in forma I per alti contenuti di esene. A concentrazioni di esene ancora più alte (18.8 mol%) la forma mesomorfa trans-planare non si ottiene neanche ad alte deformazioni, ma si osserva la trasformazione della forma I presente nei film pressofusi, sebbene con bassa cristallinità, in una nuova forma mesomorfa con

catene in conformazione elicoidale. Dato che tutti i campioni dei copolimeri sPPHe presentano proprietà elastiche, l’elasticità ha diversa origine a seconda della composizione. Conclusioni Questo studio ha evidenziato che il comportamento elastomerico dei copolimeri sindiotattici propilene-esene è il risultato di due diversi contributi, come già evidenziato per copolimeri sindiotattici propene-etilene4. Per campioni con un basso contenuto di esene ed elevata cristallinità l’elasticità è essenzialmente di natura entalpica (come nell’omopolimero), dovuta alla transizione che avviene durante il rilassamento dalla forma III metastabile alle forme elicoidali più stabili. Per campioni poco cristallini con alti contenuti di esene (>6.4mol%), il contributo entalpico all’elasticità è trascurabile, e la sua natura è puramente entropica. Il contributo entropico è essenzialmente dovuto alle transizioni conformazionali delle catene della fase amorfa, tra le conformazioni estese, che si formano nello stiro, e le conformazioni random coil (entropicamente favorite) che si riformano dopo rimozione della tensione. Infine per i campioni con un contenuto di esene intermedio (tra 3.9mol% e 5.6mol%) il comportamento elastico può essere attribuito sia al contributo entalpico che a quello entropico. Il contributo entalpico riveste un ruolo sempre meno incisivo all’aumentare della concentrazione di esene. Riferimenti 1. M. Kakugo, Macromol. Symp., 89, 545 (1995). 2. N. Naga, K. Mizunuma, H. Sadatashi, M. Kakugo, Macromolecules, 30, 2197 (1997). 3. F. Auriemma, O. Ruiz de Ballesteros, C. De Rosa, Macromolecules, 34, 4485 (2001). 4. C. De Rosa, F. Auriemma, Advanced Materials, 17, 1503 (2005).


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COPOLIMERI A BLOCCHI POLI(BUTILENE/DIETILENE SUCCINATO): EFFETTO DELLA LUNGHEZZA DEI BLOCCHI SULLA MISCIBILITA’, SUL COMPORTAMENTO DI FASE E SULLA CAPACITA’ A CRISTALLIZZARE

N. Lotti, M. Soccio, L. Finelli, A. Munari. Università degli Studi di Bologna, Dipartimento di Chimica Applicata e Scienza dei Materiali, Viale Risorgimento 2, 40136, Bologna, Italy; e-mail: [email protected].

Introduzione Gli enormi progressi ottenuti nell’ultimo decennio nel campo dei Biomateriali e della Biologia Cellulare hanno consentito di passare dalla medicina riparativa a quella rigenerativa ed alle terapie cellulari. In questo contesto lo sviluppo di materiali sintetici è diretto verso la progettazione di sistemi di trasporto, di sistemi di supporto e verso la realizzazione di materiali “intelligenti” sensibili a stimoli fisiologici. Nella presente lavoro l’attenzione è stata rivolta verso i poliesteri alifatici, notoriamente biodegradabili e ed in gran parte bioriassorbili, ottenuti a partire da glicoli ed acidi bicarbossilici alifatici. Appartiene a questa classe di polimeri il poli(butilene succinato) (PBS) che, insieme ad alcuni suoi copolimeri, si è rivelato un materiale promettente per svariati impieghi, tra cui quello come materiale termoplastico biodegradabile e quello in campo biomedico. Il notevole interesse mostrato verso questo polimero, che è commercializzato da diverse ditte, è imputabile al fatto che il PBS presenta proprietà simili al polietilene. Recentemente, alcuni Autori hanno proceduto alla sintesi e caratterizzazione di copolimeri statistici poli(butilene/dietilene succinato) (1). L’omopolimero poli(dietilene succinato) (PDGS) differisce dal PBS per la presenza di un atomo di ossigeno etereo nell’unità ripetitiva del polimero. La presenza di tale eteroatomo lungo la catena polimerica potrebbe aumentare l’idrofilicità del materiale finale rendendolo, ad esempio, più adatto per applicazioni nel campo dell’ingegneria tessutale. Come noto, la miscelazione reattiva di due polimeri riscuote, a tutt’oggi, un notevole interesse applicativo, in quanto consente di ottenere nuovi materiali, potenzialmente in numero illimitato ed a costi decisamente più contenuti, aventi proprietà mirate al tipo di utilizzo, partendo generalmente da polimeri già disponibili in commercio. In quest’ottica si è ritenuto interessante procedere alla miscelazione in fuso di PBS e PDGS, sintetizzati nei nostri laboratori, allo scopo di preparare copolimeri a blocchi con diversa lunghezza di catena dei due omopolimeri di riferimento. A tal fine, sono stati scelti un’adeguata temperatura di mescolamento ed opportuni tempi di mescolamento. I materiali ottenuti e i due omopolimeri di riferimento sono stati quindi sottoposti ad un’accurata caratterizzazione molecolare, tramite 1H-NMR e GPC, e termica, mediante TGA e DSC, allo scopo di correlare le proprietà termiche alla lunghezza dei blocchi. Sono state anche condotte misure diffrattometriche ad alto angolo per analizzare la fase cristallina presente. Risultati Sono state preparate miscele contenenti quantità equimolari di PBS e PDGS: sono state indicate PBSPDGSX, dove X indica il tempo di mescolamento, espresso in minuti. Allo scopo di ottimizzare le condizioni di reazione, preliminarmente sono state

condotte prove a diverse temperature (165, 215, 225 e 233°C), variando il tempo di reazione ed impiegando in tutti i casi come catalizzatore Ti(OBu)4. Dai risultati ottenuti si è appurato che la temperatura ottimale di mescolamento è 225°C. I campioni ottenuti sono riportati in Tabella I. Tabella I: dati di caratterizzazione molecolare.

Polimeri Mna PDGS (mol %)

(1H-NMR) LBS

LDGS b

PBS 38300 0 - - - PDGS 28200 100 - − −

PBSPDGS5 22800 50 - - - PBSPDGS15 21500 50 25 25.6 0.079 PBSPDGS30 20200 50 11 11 0.181 PBSPDG45 21500 50 4.2 4.6 0.46 PBSPDG60 23600 50 3.7 3.8 0.534 PBSPDG90 26200 50 2.5 2.5 0.803

PBSPDGS120 25200 50 2 2 0.985 a ottenuti da misure GPC Come evidenziato dai dati di Mn riportati in Tabella I, i copolimeri sono caratterizzati da pesi molecolari alti, comparabili tra loro e con quelli dei due omopolimeri di riferimento. La struttura chimica di tutti i campioni è stata determinata tramite spettroscopia 1H-NMR. Tutti i copolimeri hanno evidenziato una struttura in accordo con quella prevista e la loro composizione effettiva è risultata simile a quella di alimentazione. La spettroscopia 1H-NMR è stata utilizzata anche per ottenere informazioni sulle variazioni nella distribuzione delle sequenze butilene succinato e dietilene succinato lungo la catena polimerica in funzione del tempo di mescolamento. Come noto, informazioni sulla disposizione delle unità comonomeriche in catena possono essere dedotte dal grado di statisticità b, che assume valore uguale ad 1 per i copolimeri statistici, 2 per quelli alternati ed è compreso tra 0 e 1 per quelli a blocchi. In Tabella I sono riportati i valori di b ottenuti per tutti i campioni in esame. Esso risulta compreso tra 0 e 1, e ciò indica la natura a blocchi dei copolimeri in esame, ad eccezione del campione PBSPDGS120 che risulta statistico. Per quanto concerne la lunghezza dei blocchi, dai valori riportati in Tabella I si evince, come prevedibile, che il prolungamento del tempo di mescolamento comporta una diminuzione della loro lunghezza (LBS e LDGS) ascrivibile al procedere delle reazioni di scambio, con una progressiva evoluzione della distribuzione delle sequenze da una struttura a blocchi a quella statistica (PBSPDGS120), in accordo con l’andamento del coefficiente di randomizzazione b.


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La caratterizzazione dei copolimeri è proseguita sottoponendo i campioni ad analisi termogravimetrica e calorimetrica. Tabella II: dati di caratterizzazione termica

1st scan 2nd scan

Polimeri Tg (°C)

∆cp (J/°C·g)

Tm (°C)

∆Hm (J/g)

Tg (°C)

∆cp (J/°C·g)

Tc (°C)

∆Hc (J/g)

Tm (°C)

∆Hm (J/g)

PBS -33.9 0.023 113.1 63 -32.4 0.104 − − 112.7 58 PDGS -25.3 0.867 − − -21.5 0.804 − − − −

PBSPDGS5 -25.4 0.450 109.9 32.4 -23.5 0.405 − − 110.3 30.4 PBSPDGS15 -25.0 0.494 110.2 34.7 -23.2 0.533 - - 110.5 30.1 PBSPDGS30 -26.8 0.505 105.1 35 -23.5 0.481 5.0 3.3 105.2 30.5 PBSPDGS45 -25.8 0.461 95.5 36 -28.4 0.666 15.1 33.1 94.7 36.5 PBSPDGS60 -28.7 0.531 81.7 32 -29.6 0.683 23.5 31.2 81.7 31.2 PBSPDGS90 -28.4 0.453 64.5 32.2 -28.5 0.698 − − − − PBSPDGS120 -28.4 0.497 46.9 27.5 -28.6 0.622 − − − −

Dalle curve termogravimetriche in aria si osserva come in tutti i campioni la perdita di peso sia del 100% ed avvenga praticamente in un unico stadio. Il PDGS risulta meno stabile termicamente del PBS, come normalmente osservato nei polimeri contenenti atomi di ossigeno etereo in catena. Il comportamento dei copolimeri è sostanzialmente simile a quello del PDGS (la stabilità termica è buona fino a circa 352°C). Dato che il processo di fusione di un polimero è influenzato dalla sua precedente storia termica, prima dell’analisi i campioni sono stati mantenuti per tre mesi a temperatura ambiente. I dati calorimetrici relativi sono riportati in Tabella II. Per quanto concerne i due omopolimeri di riferimento, il PBS risulta un materiale semicristallino, come evidenziato dalla relativa curva calorimetrica, caratterizzata da una variazione endoterma della linea di base associata alla transizione vetrosa e da una cospicua endoterma di fusione. Il PDGS, al contrario, è completamente amorfo. I copolimeri oggetto del presente studio appaiono semicristallini. Per quanto concerne il fenomeno della fusione, la posizione del picco risulta dipendere dal tempo di mescolamento e quindi dalla lunghezza dei blocchi. In particolare, la temperatura di fusione diminuisce regolarmente all’aumentare del tempo di mescolamento, ovvero al diminuire della lunghezza dei blocchi, in accordo con la formazione di una fase cristallina caratterizzata da un minor grado di perfezione. Misure diffrattometriche hanno permesso di evidenziare che la fase cristallina presente nei copolimeri è quella del PBS. Per quanto riguarda il fenomeno della transizione vetrosa, in tutti i casi si individua un’unica variazione endoterma della linea di base e ciò indica che il sistema copolimerico PBSPDGS è omogeneo nella fase amorfa, anche se non completamente miscibile, essendo caratterizzato da una separazione di fase, con formazione di domini cristallini. Considerazioni più certe sullo stato di miscibilità del sistema nella fase amorfa sono state effettuate sui campioni raffreddati rapidamente dallo stato fuso. Il PBS ed i copolimeri ottenuti per tempi di mescolamento inferiori o pari a 45 minuti risultano tutti semicristallini; PBSPDGS60, PBSPDGS90 e PBSPDGS120 sono al contrario amorfi. Come mostrato in Tabella II, in tutti i casi è stata evidenziata una sola Tg a riprova della totale miscibilità nello stato amorfo del sistema PBSPDGS. Come si può notare dai valori riportati in Tabella II i

copolimeri semicristallini sono caratterizzati da una Tg più alta di quella dei copolimeri completamente amorfi: tale risultato non sorprende se si tiene conto che la fase amorfa si arricchisce in PDGS a seguito della segregazione per cristallizzazione del PBS. Le curve calorimetriche dei copolimeri PBSPDGS30, PBSPDGS45 e PBSPDGS60 presentano oltre al fenomeno della transizione vetrosa anche un picco esotermico di cristallizzazione in scansione seguito da un picco endotermo di fusione a più alta temperatura. Come evidenziato dai dati riportati in Tabella II, il procedere della miscelazione reattiva provoca un aumento della temperatura di cristallizzazione in riscaldamento ed una diminuzione della temperatura di fusione. Infatti con la progressiva diminuzione della lunghezza dei blocchi, i cristalli di PBS si formano a temperature sempre più elevate durante la scansione e ciò indica che la velocità di cristallizzazione non isoterma dei blocchi di PBS diminuisce. Questo andamento può essere spiegato sulla base di due effetti: da un lato un innalzamento della Tg e un abbassamento della Tm del copolimero rispetto all’omopolimero di riferimento fanno sì che la finestra Tg/Tm si restringa, ostacolando così la cristallizzazione dei segmenti di PBS; dall’altro la diminuzione della lunghezza dei blocchi implica che i segmenti di PBS sono più bloccati dalla presenza dei blocchi di PDGS e quindi hanno maggiore difficoltà a riarrangiarsi in un assetto ordinato. Nei copolimeri PBSPDGS90 e PBSPDGS120 i blocchi diventano così corti che nelle condizioni sperimentali adottate i segmenti di PBS non hanno la capacità di formare cristalli. Conclusioni I risultati più salienti della ricerca possono essere così riassunti: � Ottimizzazione del processo di produzione del materiale con la possibilità di “pilotare” la sintesi e ottenere copolimeri con struttura a blocchi o statistica. � La stabilità termica, proprietà cruciale durante la fase di lavorazione di un polimero, si mantiene buona per tutti i campioni preparati. � Il sistema è omogeneo nella fase amorfa, anche se non completamente miscibile, essendo caratterizzato da una separazione di fase, con formazione di domini cristallini di PBS. � La capacità a cristallizzare in riscaldamento a partire dalla fase amorfa dei segmenti di PBS è strettamente correlata alla lunghezza dei blocchi e quindi al tempo di mescolamento; al di sotto di un certo valore critico i blocchi di PBS non riescono a riarrangiarsi in un assetto ordinato. Riferimenti A. Cao, T. Okamura, K. Nakayama, Y. Inoue, T. Masuda, Polymer Degradation and Stability, Vol 78, 107-117 (2002).


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NEW COPOLYMERS OF ω−PENTADECALACTONE WITH P-DIOXANONE BY LIPASE CATALYSIS

M.L. Focarete1, M. Scandola1, Z. Jiang,2 H. Azim2, R.A. Gross2 1 Università di Bologna, Dipartimento di Chimica “G.Ciamician”, via Selmi 2, 40126 Bologna, Italy; e-mail: [email protected]

2 Polytechnic University, Department of Chemistry and Chemical Engineering, Six Metrotech Center, Brooklyn, NY-11201, USA

Introduction Copolymerization is an effective strategy to obtain materials with improved properties compared with those of the respective homopolymers. Poly(p-dioxanone) (PDO), and its copolymers with caprolactone, glycolide, lactide, are well known bioresorbable materials. Current processes to synthesize PDO and DO-containing copolyesters make use of organometallic catalysts. However, in order to obtain medical-grade polymers, residual metal content in the polymer should be minimized or completely eliminated. Metal-free enzymatic catalysis is a valuable method for preparing polymers for medical applications. Lipase catalysis is known to be highly efficient in several lactone ring-opening polymerizations.1 Copolymers of DO with large ring size lactone monomers, such as ω-pentadecalactone (PDL) have not been reported so far. In this work the synthesis of new P(PDL-co-DO) copolymers, as well as their thermal and crystalline properties are reported. Results High molecular weight poly(PDL-co-DO) (weight average molar mass > 30 kDa) were obtained through the enzymatic copolymerization of PDL and DO monomers using Candida antarctica Lipase B. Proton (1H) and carbon-13 (13C) NMR analyses showed that, due to transesterification promoted by the lipase, copolymers containing nearly random sequences of PDL and DO units were obtained. Solid-state properties of poly(PDL-co-DO), as well as those of the reference homopolymers, were investigated by thermogravimetry (TGA), differential scanning calorimetry (DSC), and X-ray diffraction (WAXS). Thermal stability of the two homopolymers is rather different, the temperature of maximum degradation rate of PPDL (Tmax = 414°C) being about 100°C higher than that of PDO (Tmax = 309°C). All copolymers, up to a content of DO units of 70 mol%, show a single weight loss degradation process, at a Tmax that decreases with increasing content of DO units. The TGA results demonstrate that copolymerization with PDL is an effective way to enhance PDO thermal stability, when the PDL content is ≥ 30 mol%. Both PPDL and PDO are semicrystalline polymers with similar degree of crystallinity, estimated to be around 60% by WAXS. The calorimetric behaviour of the two homopolymers is different due to differences in the crystallization kinetics. PPDL is a polymer that crystallizes so quickly from the melt that it cannot be quenched effectively enough to give a sample with a

large amorphous fraction. As a consequence, the glass transition temperature (Tg) of PPDL is not easily obtained by DSC but is better revealed by relaxation techniques (Tg = - 27°C by dynamic mechanical analysis).2 On the contrary, PDO can be easily quenched from the melt and it shows an intense glass transition at Tg = - 10°C (by DSC). Both homopolymers are characterized by a melting endotherm at a similar temperature (Tm PPDL = 93°C, Tm PDO = 107°C), but with different melting enthalpies: ∆Hm PPDL = 127 J/g, ∆Hm PDO = 82 J/g. All poly(PDL-co-DO) are crystalline over the whole composition range (crystallinity degree in the range from 44 % to 59% by WAXS). They crystallize in either of the crystal lattices of the respective homopolymers. Depending on copolymer composition, either PDO or PPDL crystal lattice is stable and is able to host units of the second monomer. For the copolymer containing around 70 mol % DO, the crystal phase changes from one lattice to the other, with both crystal types found side by side. This particular copolymer composition represents the pseudoeutectic, where both melting temperature and melting enthalpy show a minimum. This behaviour suggests that poly(PDL-co-DO) is an isodimorphic system. Conclusions In this work, new copolymers of i-pentadecalactone and p-dioxanone were successfully synthesized by lipase catalysis under mild reaction conditions. Poly(PDL-co-DO) copolymers show isodimorphic behaviour and are partially crystalline over the whole range of composition. Given the difference in hydrofilicity of the comonomers (PDL and DO), the hydrolytic degradation rate of the copolymers can be modulated by changes in copolymer composition, producing materials for targeted biomedical applications. References 1. R.A. Gross, A. Kumar, B. Kalra Chem. Rev. 101,

2097-2124 (2001) 2. M.L. Focarete, M. Scandola, A. Kumar, R.A. Gross

J. Polym.Sci., Part B: Polym. Phys., 39, 1721-1729. (2001)


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SYNTHESIS AND CHARACTERIZATION OF POLYBENZOFULVENE DERIVATIVES ROLE OF SUBSTITUENTS IN THE MODULATION OF POLYMER PROPERTIES

R. Mendichi 1, L. Zetta 1, A. Cappelli 2, S. Galeazzi 2, G. Giuliani 2, M. Anzini 2, A. Donati 3 1 Istituto per lo Studio delle Macromolecole (CNR), Via E. Bassini 15, 20133 Milano, e-mail: [email protected]

2 Dipartimento Farmaco Chimico Tecnologico, Università di Siena, Via A. Moro, 53100 Siena 3 Dipartimento Scienze Chimiche e Biosistemi, Università di Siena, Via A. Moro, 53100 Siena

Introduction In the chemistry of the advanced materials the synthesis of vinyl polymers possessing stable i-stacked conformation represents an important objective and in this context the properties of polydibenzofulvene (PDBF) was investigated. We recently reported [1, 2] that a benzofulvene derivative [ethyl 1-methylene-3-(4-methylphenyl)-1H-indene-2-carboxylate undergo sponta-neous polymerization to give poly-BF1, see Scheme 1, a polymer showing interesting properties such as thermo-reversible polymerization/depolymerization behaviour, high solubility in most organic solvents, susceptibility to molecular manipulation and characterized by both vinyl structure stabilized by aromatic stacking interactions and very high molar mass. This presentation describe more important results of a study focused on the macromole-cular structure and properties together to the polymeri-zation mechanism of some PDBF derivatives. Results A general method for the synthesis of benzofulvene monomers is reported in Scheme 1 [2, 3]. In particular poly-BF1 shows interesting macromolecular properties.

Scheme 1 Polymerization/depolymerization behaviour of poly-BF1 polymer was described in detail previously [1, 2]. The poly-BF1 polymerizes spontaneously upon removal of solvent and very interesting shows thermo-reversible polymerization. Because the spontaneous polymerization of BF1 occurs upon the removal of the solvent, the polymerization conditions are characterized by an instantaneous change in the monomer concentration. Under these conditions it is reasonable to assume that the monomer concentration affects both the initiation and propagation steps. The molecular characterization of freshly prepared poly-BF1 (different batches, same methodology but different amounts of monomer) was performed with a multi-angle laser light scattering detector on-line to a size exclusion

chromatography system (SEC–MALS). The Mw average of four different poly-BF1 samples ranged from ultrahigh (2670 kg/mol) to high (867 kg/mol) molar mass, depending on the amount of the monomer used in the polymerization (Table 1). Table 1. SEC-MALS results of four poly-BF1 samples. Sample Mw Mw/Mn Rg b Kg/mol nm P100 a 867 2.0 30.8 P200 a 1,605 1.8 43.1 P400 a 2,670 1.6 60.2 P800 a 1,437 1.8 41.0 MALDI-TOF spectrum of poly-BF1 confirmed the polymeric nature of this material. The MALDI spectrum (Figure 2) shows the presence of a series of peaks differing by 290 (molar mass of the monomer BF1) up to a m/z value of about 8000. Heavier polymeric species were undetectable under MALDI-TOF conditions.

Figure 2. MALDI-TOF of poly-BF1.

0.1

1.0

10.0

1.0E+05 1.0E+06 1.0E+07 1.0E+08Molar Mass [g/mol]

[η]dL/g

10

100

1000Rgnm[η]=ƒ(M)

a = 0.85

Rg=ƒ(M)α = 0.63

Figure 3. MHS plot and conformation plot of POLY-BF1 by a SEC-MALS-DV system.


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The conformation of poly-BF1 polymer is stiff, see Figure 3, as demonstrated by the slope of the MHS plot, [i]=ƒ(M), and the conformation plot, Rg=ƒ(M), obtained by a SEC-MALS-DV system (where DV denotes an on-line differential viscometer). Thermal analysis, both TGA [2] and DSC (see Figure 4), shows that poly-BF1 in the solid state is not subject to thermal degradation up to approximately 210 °C (weight loss lower than 2%). The fast thermal degradation of the poly-BF1 starts approximately at 240 °C with a maximum of the derivative at about 250 °C.

Figure 4. DSC curve of poly-BF1. A viable method for the synthesis of novel benzofulvene derivatives bearing different substituents in key positions of indene moiety is reported in Scheme 2.

Scheme 2 A summary of the macromolecular features of some synthesized PDBF derivatives are shown in Table 2. The structure of the newly synthesized PDBF derivatives was investigated with NMR [1-3]. The simplest situation was found for poly-3k, which resulted to be very similar to poly-BF1 (i.e. poly-3l in Table 2). 13C NMR spectra of poly-3k,l showed a single peak attributable to the aliphatic quaternary carbon with a chemical shift value (about 58 ppm) close to both the one reported for poly-BF2 (57.8 ppm) and the one calculated for poly-BF1 in the 1,2-enchainment (54.2 ppm. Moreover, the broad peak centered at 48.4 ppm attributed to the methylene bridge of poly-BF1 was resolved in the spectra of both polymers (poly-3k,l) into two broad peaks (at about 48 and 51 ppm), showing different intensity and suggesting the contemporary presence of a minor species with a major one. The comparison of the chemical shifts of these two broad peaks with that found in the case of poly-BF2 (43.8 ppm) and with those calculated for the 1,2-enchainment (45.6 ppm) and the 1,4-enchainment (30.2 ppm) of poly-BF1 supports the vinyl nature (1,2-enchainment) of poly-3k,l and suggests the possible existence of configurational or conformational isomers. When the bulky electron-withdrawing COOC2H5 of poly- 3k,l was replaced by a small donor methyl group in

poly-3f, the 13C signals attributable to the carbon atom of the backbone methylene and to the aliphatic quaternary carbon in position 1 of indene moiety of poly-3f were clearly doubled. It is noteworthy that the chemical shift values of the couple of signals attributable to the aliphatic quaternary carbon of poly-3f (57.4 and 61.0 ppm) were very close to those calculated for poly-BF1 in the two different enchainments (54.2 and 63.1 ppm). Thus, in some monomeric units of poly-3f the indene double bond could be retained in the original position, while in some others it could be shifted as a consequence of a diene (1,4) polymerization. On the other hand, several minor species are present in the 13C NMR spectra of poly-3g,h and poly-3c-e. The number of the peaks attributable to aliphatic quaternary carbon and their chemical shift suggest the presence in these polymers of enchainment, configurational, and/or conformational isomers. Finally, these results suggest that the steric hindrance in the proximity of the reactive exocyclic methylene group of the monomer affects the macromolecular structure of these polymers. Table 2. Description of some PDBF derivatives. Polymer R1 R2 R3 Mw D Rg poly-3a H H H 4.0 1.3 poly-3c F H H 25.6 4.5

poly-3d Cl H H 47.6 2.1 poly-3e Br H H 109.1 3.1 poly-3f CH3 H H 530.8 4.7 43.0 poly-3g CN H H 391.5 6.3 32.5 poly-3h CN CH3 H 500.2 4.4 26.8 poly-3i CN H CH3 475.9 4.3 24.3 poly-3j C≡C-2-Pyr H H 289.2 2.5 poly-3k COOC2H5 H H 1,593.0 3.7 44.6 poly-3l COOC2H5 CH3 H 2,670.0 1.6 poly-3ma COOH CH3 H 694.9 9.6 Conclusions Synthesis, structure and macromolecular properties of some PDBF derivatives have been investigated by light scattering, viscometry, NMR, MALDI-TOF and thermal analysis. In particular poly-BF1 shows interesting properties such as thermo-reversible polymerization/de-polymerization behaviour, high solubility in organic solvents, susceptibility to molecular manipulation. References [1] A. Cappelli, G. Pericot Mohr, M. Anzini, S. Vomero, A. Donati, M. Casolaro, R. Mendichi, G. Giorgi, J. Org. Chem., 68, 9473 (2003) [2] A. Cappelli, M. Anzini, S. Vomero, A. Donati, L. Zetta, R. Mendichi, M. Casolaro, P. Lupetti, P. Salvatici, G. Giorgi, J. Polym. Scie. A, 43, 3289 (2005). A. Cappelli, S. Galeazzi, G. Giuliani, M. Anzini, A. Donati, L. Zetta, R. Mendichi, M. Aggravi, G. Giorgi, E. Paccagnini, S. Vomero, Macromolecules, 40, 3005 (2007).

XVIII Convegno Italiano di Scienza e Tecnologia delle Macromolecole, Catania 16-20 settembre 2007 S2 P6/18

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CARATTERIZZAZIONE CHIMICO-FISICA DI POLIFENILENI ETINILENI CRISTALLINI E LIQUIDO CRISTALLINI

L. D’Ilario, A. Martinelli, A. Micozzi Università di Roma “La Sapienza”, Dipartimento di Chimica, P.le A. Moro, 5, 00185, Roma, Italy, e-mail:

[email protected]

Introduzione Sistemi coniugati polimerici costituiti dall’alternanza di alchini ed arileni, i poliarileni etinileni (PAE), rappresentano una classe di polimeri che, in virtù dell’estesa coniugazione del sistema π, mostrano proprietà peculiari quali una piccola bandgap energy, alta rigidità e facilità di trasporto di carica. Queste caratteristiche conferiscono a tali materiali interessanti proprietà optoelettroniche di foto- ed elettro-luminescenza, che possono essere opportunamente impiegate per la realizzazione di dispositivi quali diodi ad emissione di luce (OLED), field-effect transistors (FET), sensori e celle fotovoltaiche. Lo spettro di emissione dei PAE può essere opportunamente modulato mediante la scelta della componente aromatica o l’incorporazione di eteroatomi o metalli nella catena coniugata mentre le caratteristiche di processabilità possono essere variate mediante l’innesto di catene alchiliche lungo lo scheletro rigido del polimero, che ne aumenta la solubilità nei solventi organici. La foto-fisica di tali polimeri e le proprietà di emissione dipendono inoltre dall’ordine a lungo raggio che si può creare tra le catene rigide e dalle loro mutue interazioni. Recentemente è stato sviluppata una strategia innovativa per la sintesi di PAE, denominata Extended One Pot (EOP)[1], che consente di preparare PAE mediante la formazione in-situ di derivati organostannanici difunzionalizzati ed il loro successivo accoppiamento diretto con dialogenuri organici in presenza di un catalizzatore di Pd. Mediante tale procedura sono stati sintetizzati tre poli(2,5-dialcossifenileni-1,4-etinileni) con le catene laterali alcossiliche di 4, 8 e 16 atomi di carbonio, denominati PPE4, PPE8, PPE16, rispettivamente. Tramite la caratterizzazione termica (DSC), ottica (microscopio ottico in luce polarizzata, POM), spettroscopica (FT-IR) e strutturale (diffrazione di raggi X in dispersione di energia ad alto e medio angolo, WA-EDXD e MA-EDXD), per il polimero PPE16 è stata ipotizzata la formazione di una fase sanidica termotropica, nella quale le catene polimeriche rigide ed i gruppi C16 lateralmente estesi, sono orientati su piani sovrapposti. Risultati e Discussione L’analisi calorimetrica ha mostrato che i campioni PPE4, PPE8 e PPE16 presentano un processo endotermico alla temperatura di circa 70, 60 e 50 °C, rispettivamente, attribuibile alla distruzione della struttura regolare delle catene laterali. Il PPE16 presenta inoltre altri due picchi endotermici a 80 e 110°C circa. Seguendo l’evoluzione morfologica dei campioni mediante l’impiego dell’analisi POM e dalla misura dell’intensità luminosa

trasmessa in funzione della temperatura si è potuto osservare che il polimero subisce un processo di fusione intorno a 80 °C. A tale temperatura il campione mantiene la birifrangenza tipica di una fase liquido-cristallina. La scomparsa di figure di birifrangenza avviene a circa 110 °C, in corrispondenza della temperatura di isotropizzazione. L’interpretazione dei fenomeni associati alle transizioni termiche è stata confermata da studi strutturali. L’analisi degli spettri di diffrazione di raggi X condotta su tutti e tre i polimeri ha evidenziato la presenza di una figura di diffrazione ad alto angolo, indice di un ordine cristallino a corto raggio, e di un picco di diffrazione ad angoli maggiori, corrispondente alla distanza intercatena. Tale distanza aumenta all’aumentare della lunghezza delle catene laterali. L’evoluzione della struttura in conseguenza di trattamenti termici è stata seguita conducendo esperimenti EDXD sul campione PPE16 alle temperature di 58, 92, 116 °C, ossia alcuni gradi al di sopra delle transizioni termiche. In generale si osserva che il picco a medio angolo si allarga e si sposta a vettori di scattering maggiori, probabilmente a causa del processo di disorganizzazione delle catene laterali che conduce all’avvicinamento delle catene principali. Esso scompare solo alla temperatura di isotropizzazione (116 °C). Rispetto al comportamento registrato a temperatura ambiente, a 58 °C la figura di diffrazione ad alto angolo permane pur subendo modifiche sostanziali, mentre si trasforma in un alone amorfo a 92 °C, ossia al di sopra della temperatura di fusione. La permanenza a tale temperatura del picco a medio angolo mette in evidenza la formazione di una fase liquido-cristallina. Combinando i dati di diffrazione, relativi alle due regioni angolari, in via preliminare si può ipotizzare per il campione PPE16 una struttura layered di tipo sanidico nella quale le catene polimeriche rigide e le catene esadeciliche, supposte lateralmente estese, sono orientate su piani sovrapposti. Seguendo la variazione dell’intensità degli assorbimenti IR del PPE16 in funzione della temperatura si sono osservate discontinuità in corrispondenza delle transizioni termiche a circa 50 e 110 °C. Solo il picco a 800 cm-1, non ancora assegnato, mostra una variazione netta di intensità a circa 80 °C. Bibliografia 1. R. Pizzoferrato, M. Berliocchi, A Di Carlo, P. Lugli, M. Venazzi, A. Micozzi, A. Ricci, C. Lo Sterzo, Macromolecules, 36, 2215 (2003)


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SINTESI DI POLI(IDROSSIALCHILENOSSI BENZOATO) A DIVERSA LUNGHEZZA DI CATENA: EFFETTO DELLA STRUTTURA CHIMICA SULLE PROPRIETA’ TERMICHE

V. Siracusa1, N. Lotti2, A. Munari2, L. Finelli2

1 Università degli Studi di Catania, Facoltà di Ingegneria, Dipartimento di Metodologie Fisiche e Chimiche per l’Ingegneria, Viale Andrea Doria 6, 95100, Catania, Italy; e-mail: [email protected]

2 Università degli Studi di Bologna, Facoltà di Ingegneria, Dipartimento di Chimica Applicata e Scienza dei Materiali, Viale Risorgimento 2, 40138, Bologna, Italy; e-mail: [email protected]

Introduzione Il poli(4-idrossietossi benzoato) (PHEBA) è un polimero a catena piuttosto rigida, parzialmente cristallino a temperatura ambiente e caratterizzato da un’elevata temperatura di fusione (1). In passato è stato usato come fibra prodotta in Giappone con il nome commerciale A-Tell. Dopo averlo sintetizzato e caratterizzato da un punto di vista chimico e termico, abbiamo ritenuto interessante proseguire l’attività di ricerca studiando l’influenza della lunghezza della catena alifatica sulle proprietà del materiale finale. In quest’ottica, preliminarmente si è proceduto alla sintesi dei monomeri contenenti tre, quattro e sei gruppi CH2 nella catena alifatica e, dopo averli purificati e caratterizzati, si è effettuata la sintesi dei rispettivi polimeri. Questi ultimi stati sottoposti ad un’accurata caratterizzazione molecolare e termica, allo scopo di valutare come l’introduzione di un numero sempre maggiore di gruppi CH2 in catena possa influenzare le proprietà dei materiali ottenuti rispetto al polimero di riferimento, il PHEBA. Risultati I monomeri sono stati sintetizzati facendo reagire l’acido p-(4-idrossibenzoico) con il 3-Cloro-propanolo (sintesi a), con il 3-Cloro-1 Butanolo (sintesi b) e con il 6-Cloro-1-Esanolo (sintesi c). I monomeri ottenuti sono stati cristallizzati più volte nel relativo solvente e quindi caratterizzati mediante indagine spettroscopica 1H.NMR, analisi elementare, analisi spettroscopica FT-IR e punto di fusione. Sintesi (a): acido p(4-idrossipropossi)benzoico (HPBA): in 225 ml di una soluzione acquosa di NaOH( 27.45 g), vengono sciolti 45 g di acido p-(4-idrossibenzoico). Successivamente a tale soluzione vengono aggiunti, goccia a goccia, 34.03 g di 3-Cloro propanolo a 50°C, per circa 45 minuti. Alla miscela di reazione, dopo averla lasciata sotto agitazione per 6 ore a 75°C, vengono aggiunti, sempre goccia a goccia e sotto continua agitazione, 15.75 ml di H2SO4 concentrato ottenendo, dopo filtrazione e lavaggio in acqua, un precipitato bianco polveroso con una resa dell’86%. Dopo due cristallizzazioni in Acetonitrile (CH3CN), si è ottenuto un composto cristallino (resa finale del 75%) che è stato pienamente caratterizzato. Sintesi (b): acido p(4-idrossibutossi)benzoico (HBBA): le condizioni di sintesi sono analoghe a quelle della sintesi a aggiungendo però 38.79 g di 3-Cloro-1 Butanolo come reagente di reazione. Dopo filtrazione e lavaggio in acqua, si ottiene un precipitato bianco con una resa dell’52%. Dopo due cristallizzazioni in CH3CN,

si è ottenuto un composto cristallino (resa finale del 35%) che è stato pienamente caratterizzato. Per aumentare la resa di reazione è stato aumentato il tempo di reazione del primo step (da 6 a 12 ore) e sono stati variati i rapporti stechiometrici acido p-(4-idrossibenzoico:NaOH (da 1:2 a 1:2.5) e acido p-(4-idrossibenzoico):3-Cloro-1-Butanolo (da1:1,1 a 1:1,5). Sintesi (c): acido p(4-idrossiesossi)benzoico (HEsaBA): come la sintesi a aggiungendo però 16.23 g di 6-Cloro-1-esanolo come reagente di reazione. Dopo filtrazione e lavaggio in acqua, si ottiene un precipitato bianco con una resa dell’82%. Dopo due cristallizzazioni in Acetone, si è ottenuto un composto cristallino (resa finale del 78%) che è stato pienamente caratterizzato. La reazione generale di sintesi è la seguente:

COOH

OH

COOH

O CH2 OH

Cl CH2 OHn

n con n=3 (HPBA), 4 (HBBA), 6 (HEsaBA) I risultati della caratterizzazione molecolare sono riportati in Tabella I. Tabella I: risultati dell’analisi elementare Campione %C

sperim. %H sperim.

%C teorico

%H teorico

HPBA 60.49 6.39 61.22 6.16 HBBA 62.54 6.66 62.55 6.68 HEsaBA 65.12 7.61 65.53 7.61

I polimeri contenenti un diverso numero di gruppi CH2 nella catena alifatica, sono stati sintetizzati in massa, impiegando l’usuale procedura a due stadi, partendo dal rispettivo monomero HPBA, HBBA o HEsaBA. L’impiego del Ti(OBu)4 come catalizzatore e le elevate temperature di polimerizzazione (220-240°C) hanno portato ad ottenere polimeri statistici per quanto riguarda la distribuzione dei pesi molecolari. I campioni sintetizzati sono stati sottoposti a caratterizzazione molecolare (analisi spettroscopica 1H.n.m.r. e cromatografia di permeazione su gel, 1100 Hewlett Packard) e termica (termobilancia Perkin ElmerTGA7 e calorimetro differenziale a scansione Perkin Elmer DSC7).


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L’unità ripetitiva dei polimeri è riportata qui di seguito:

C OO

O

CH2 n

m n=3, P-HPBA n=4, P-HBBA n=6, P-HEsaBA I dati di caratterizzazione molecolare (Mn) e dell’analisi termogravimetrica (TGA) sono riportati in Tabella II mentre nella Tabella III sono riportati i dati di caratterizzazione termica ottenuti mediante analisi DSC. Tabella II: dati di caratterizzazione molecolare

Polimeri Mn

Tid (°C)

Tmax (°C) PHEBA 9500 412 446 PHPBA 9200 390 415 PHBBA 8600 389 421

PHEsaBA 10100 393 417 Tabella III: dati di caratterizzazione termica

Dall’analisi termogravimetrica è emerso che tutti i polimeri presentano una buona stabilità termica (fino a 390°C) analogamente all’omopolimero di riferimento PHEBA, benchè vi sia un leggero effetto dovuto all’allungamento della catena. Infatti, la stabilità termica degli omopolimeri sintetizzati è risultata dipendere, seppure in modo lieve, dalla lunghezza della catena alifatica, essendo tanto minore quanto maggiore è il numero di gruppi CH2 presenti per unità ripetitiva (vedi Figura 1). Dall’analisi delle curve calorimetriche (vedi Figura 2) e dai dati riportati in Tabella II si nota che il valore della Tg diminuisce all’aumentare del numero dei gruppi CH2 in catena: l’effetto è imputabile ad un aumento della flessibilità della catena polimerica. Per quanto concerne il fenomeno della fusione, il valore di Tm risulta dipendere dalla struttura chimica: si può notare un effetto pari-dispari. Relativamente ai campioni contenenti un numero pari di gruppi metilenici, si osserva una diminuzione di Tm all’aumentare della lunghezza di catena, imputabile ad un aumento di flessibilità della catena stessa. La più elevata temperatura di fusione osservata per PHPBA è probabilemente da ricondurre alla particolare struttura cristallina. Sono in corso indagini mediante diffrattometria a Raggi X per convalidare tale ipotesi.

50 150 250 350 450 550 650 750T (°C)

0

20

40

60

80

100

WEI

GHT (

%)

PHEBAPHPBAPHBBAPHesaBA

Figura 1: curve termo gravimetriche in aria, 10°C/min.

0 50 100 150 200T (°C)

endo

→

PHEBA

PHPBA

PHBBA

PHesaBA

Figura 2: curve calorimetriche dopo raffreddamento veloce dallo stato fuso (20°C/min). Dai dati riportati in Tabella III si evince che anche il comportamento di fase risente della struttura chimica: dopo rapido raffreddamento dallo stato fuso, infatti, PHEBA e PHBBA risultano completamente amorfi, anche se, una volta superata Tg le catene macromolecolari hanno sufficiente mobilità ed energia per riarrangiarsi e cristallizzare, mentre PHPBA E PHEsaBA sono parzialmente cristallini. L’elevata velocità di cristallizzazione del PHPBA può essere ricondotta alla particolare struttura cristallina; l’aumentata capacità a cristallizzare del PHEsaBA rispetto a PHEBA e PHBBA è invece imputabile alla maggiore flessibilità di catena del polimero a catena alifatica più lunga. Conclusioni I risultati ottenuti mostrano come la struttura chimica sia un parametro chiave nella modulazione delle proprietà termiche del materiale, in particolare: � stabilità termica; � temperatura delle transizioni termiche caratteristiche; � comportamento di fase; � velocità di cristallizzazione. Riferimenti 1. N. Lotti, V. Siracusa, L. Finelli, A. Munari, P.

Manaresi, Journal of Thermal Analysis and Calorimetry, Vol. 66, 827-840 (2001).

Polimeri Tg (°C)

∆cp (J/°C·g)

Tc (°C) ∆Hc

(J/g) Tm

(°C) ∆Hm (J/g)

PHEBA 75 0.47 129 48 209 49 PHPBA 59 0.13 160 14 210 57 PHBBA 37 0.48 84 43 156 45 PHEsaBAA

18 0.23 31 26 145 49


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THERMAL AND STRUCTURAL INVESTIGATION OF TRANS-1,4 BUTADIENE-ISOPRENE COPOLYMERS

F. Bertini, M. Canetti, G. Ricci

CNR-Istituto per lo Studio delle Macromolecole (ISMAC), Via E. Bassini 15, 20133 Milano, Italy

e-mail: [email protected]

Introduction Trans-1,4 polybutadiene (PB) is at room temperature a hard, crystalline material; its high melting temperature has a negative effect on its processing properties. Nevertheless, it is known that high trans-1,4 diene polymers present good dynamic properties. With the aim to reduce the crystallinity and melting temperature of PB, isoprene units were introduced in the polymer backbone [1]. Butadiene-isoprene copolymers with trans-1,4 structure were prepared at –30°C in the whole range of monomeric ratio, using the catalyst system V(acac)3-MAO (Table I). In the present paper, we report the results of a detailed thermal and structural investigation performed by DSC, TGA and WAXD. Table I: Copolymers composition and characteristics Sample code XBa) Mw ·10-3 b) Tg (°C)

PB 1 510 - C1 0.98 500 -82.1 C2 0.96 495 -81.9 C3 0.90 480 -80.6 C4 0.85 450 -80.2 C5 0.83 440 -81.4 C6 0.77 355 -80.4 C7 0.75 365 -80.8 C8 0.67 370 -78.6 C9 0.54 370 -78.2 C10 0.46 350 -77.3 C11 0.20 280 -72.6 PI 0 136 -68.0

a) Butadiene fraction determined by 13C NMR. b) Determined by GPC analysis.

Results The samples were cooled from the melt to -120°C, then heated to 150°C at constant rate of 20°C/min. In Figure 1 the DSC heating curves of isoprene and butadiene homopolymers, and of selected butadiene-isoprene copolymers are shown. Polyisoprene (PI) sample presents a single melting event centred at 38°C, whereas two endothermic peaks due to the monoclinic-hexagonal transition and to the melting, respectively, are observed for pure PB. It can be observed that the monoclinic-hexagonal and melting transitions in the copolymers occur at lower temperatures compared to those registered for pure PB. A progressive

decrease of the temperature interval between the two endothermic transitions takes place with the increase of the isoprene content. The copolymers with isoprene content ranging from 23 to 33 mol-%, show only one endothermic peak. Finally, no first order transition is observed for copolymers containing 46, 54, and 80 mol-% of isoprene.

Figure 1. DSC heating scans of PB, PI and some butadiene-isoprene copolymers The glass transition phenomenon was observed for pure PI; on the contrary, quenched pure PB does not show any thermal events preceding the monoclinic-hexagonal transition. The Tg values of the copolymers are influenced by composition and by monomers distribution along the polymer chain (Table I). The experimental Tg values are well fitted by Johnston equation allowing to calculate a Tg value of -82.3°C for pure PB. Investigations on the monoclinic crystal organization and the hexagonal mesophase of pure PB and copolymers, were carried out by WAXD. Figure 2 reports the diffraction profiles for PB homopolymer and some copolymers registered at 22 °C. Pure PB and C1 copolymer show the typical profile of trans-1,4 polybutadiene crystallized in the monoclinic form. DSC experiments evidenced that the monoclinic-hexagonal transition shifted to lower temperature by increasing the isoprene content in the copolymer.

-120 -70 -20 30 80 130Temperature (°C)

endo

>>

PB

C2

C5

C6

C9

PI


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Figure 2. WAXD profiles for PB and some butadiene-isoprene copolymers. The sample C2 exhibits a WAXD pattern where the monoclinic form is predominant and a small shoulder is evidenced at about 21 °2θ. For this sample the temperatures of monoclinic-hexagonal transition (Tmh) was found at 33°C, although the transition started at about 20 °C. Reasonably, this would explain the simultaneous presence of the two crystalline forms. The C3 sample, characterised by a Tmh equal to 21°C, presents a WAXD pattern where the transition to the hexagonal form is almost complete, as revealed by the development of the peak centred at 20.8 °2θ. The presence of the small residual peak at 22.4 °2θ indicates the persistence of the monoclinic form.

The WAXD spectrum of C4 copolymer shows the typical reflection of the hexagonal form and a relevant broad scattering due to the amorphous phase. In fact, this sample has to be considered partially melted at the temperature adopted for WAXD measurements. Completely amorphous structure is evidenced for the C6 sample, characterized by an endothermic peak centred at –4°C. The effect of the isoprene content on the thermal stability of the copolymers was studied by means of TGA carried out in inert atmosphere. The copolymer composition plays an important role on the thermal degradation of the polymers: the higher is the isoprene content, the lower is the thermal stability. The decomposition of all the copolymer samples takes place in two different stages: a series of cyclization and cross-linking reactions occur before the depolymerization step. The temperature range of the events and the weight loss involved strongly depend on the copolymer composition. Conclusions The structure and the thermal behaviour of the various butadiene-isoprene copolymers are strongly influenced by composition. The attitude of the butadiene sequences to crystallize in the monoclinic form and to evolve in the hexagonal form is preserved in the copolymer for a certain range of composition. The monoclinic-hexagonal transition and melting temperatures in the copolymers shift to lower values with respect to those registered for pure polybutadiene. The temperature interval between the two endothermic events is progressively reduced by increasing the isoprene content, and for a certain copolymer composition only one endothermic peak is detected. Reference 1. G. Ricci, L. Zetta, E. Alberti, T. Motta, M. Canetti, F. Bertini, J. Polym. Sci. Part A: Polym. Chem. in press

5 10 15 20 25 30 352θ °

I (a.u

.)

PB

C1

C2

C3

C4

C6


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THE DYNAMICS OF POLYMERIC THIN FILMS

A. D'Amore, L. Grassia Dip. di Ingegneria Aerospaziale e Meccanica, Seconda Università di Napoli, Via Roma 19, Aversa, 81031 (CE) -

ITALY The study of thin polymer films is a fascinating and active area of research. Many of the studies in this area are motivated by advances in nanotechnology and the result of such studies is the production of an increasing number of technologically important systems on ever decreasing lengthscale. To fully capture the physics of these systems, an understanding of the properties of materials at these lengthscales is required. It is certainly not obvious that the properties of such highly confined materials will be the same as those of the corresponding bulk material, or even predictable from the bulk properties. Of particular interest have been studies on the dynamics in thin polymer films, where quantities such as the glass transition temperature (Tg) and the coefficient of thermal expansion display significant deviation from bulk behaviour. In this work, despite that full understanding of the fundamentals of glass formation and glassy behaviour remains one of the major unsolved problems of condensed matter physic, an attempt to model the anomalous behaviour of thin polymeric films is presented. Here the thin films under concern have

thicknesses of order of magnitude of the radius of gyration of the random coil. The approach is based on the modification of a recent theory (1) developed to predict the response of mechanically stimulated glassy polymers in their bulk form. The numerical simulations show that coupling the structural and viscoelastic relaxation phenomena the anomalous thin film behaviour can be reliably predicted. In particular, comparing our findings with the available literature data in terms of glass transition decrease, a quantitative agreement is found. Concerning the deviation of the thermal expansion coefficient from bulk behaviour, our model predicts, for supported thin films, a negative coefficient of thermal expansion (NCTE) in the glassy state even if the occurrence of a NCTE was not reported unambiguously in the relevant literature data. 1) L.Grassia, A. D’Amore “Costitutive law describing

the phenomenology of subyield mechanically stimulated glasses” Physical Review E, 74021504 (2006)


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MODELLING THE YIELD STRESS AND THE POISSON’ S RATIO OF GLASSY POLYMERS

A. D’Amore1, L. Grassia1, D. Acierno2

1Department of Aerspace and Mechanical Engineering, The Second University of Naples, Via Roma 29, 81031 Aversa (CE)

2Department of Materials and Production Engineering, University of Naples Federico II, P.le Tecchio, Naples We have recently reported that the phenomenological intricacies of mechanically stimulated glasses, even in their subtle features, can be reliably predicted by coupling the structural relaxation kinetics and viscoelastic response [1]. Here, it will be shown that our approach can be extended to predict the stress-strain behaviour of glassy polymers. In particular, a linear dependence of yield stress on the logarithmic of the strain rate is obtained, conforming to the classical Eyring-type behaviour. Furthermore the distribution of Poisson’s ratio delay times in comparison with the distribution of retardation times in shear and/or in bulk is reported, a case never appeared before. Simulations are obtained linking the modified KAHR equation with the constitutive equation for linear viscoelastic materials written in the reduced time domain.

The model requires the experimental determination of only one linear viscoelastic function (namely, the master curve of the shear relaxation obtained above Tg). Concerning the volumetric part of the viscoelastic response, it will be shown that, in the reduced time domain, the dimensionless bulk compliance coincides with the memory function appearing in the modified KAHR equation, so that it can be extracted directly from the experimental PVT data. Therefore the timescales of the responses of mechanically stimulated glasses are already encrypted into the PVT and the shear relaxation behavior. [1] L. Grassia and A. D’Amore, “Constitutive law describing the phenomenology of subyield mechanically stimulated glasses”, Physical Review E, 74, 021504 (2006)


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EFFECT ON STRUCTURAL RELAXATION OF THE PMMA COPOLYMERS CHAIN FLEXIBILITY

S. Della Sciucca1, G. Spagnoli1, M. Penco1, L. Sartore1, F. Samperi2, S. D’Antone3

1 Dipartimento di Chimica e Fisica per l’Ingegneria e per i Materiali, Università degli Studi di Brescia,Via Valotti 9, 25133 Brescia, Italy.

2 Istituto di Chimica e Tecnologia dei Polimeri (CNR), Viale A. Doria 6, 95125 Catania, Italy. 3 Dipartimento di Chimica e Chimica Industriale, Università degli Studi di Pisa, Via Risorgimento 35, 56100 Pisa,

Italy

Intoduction The structural relaxation of polymers depends on the kinetic character of the glass-transition phenomenon: amorphous polymers below their Tg are not at equilibrium and their structures continuously relax in attempt to reach the equilibrium state. This structural recovery of the homogeneous materials is obviously related to the macromolecular architecture and molecular weight distribution. In this work the enthalpy relaxation of PMMA, PMMA-co-polyethylmetacrylate and PMMA-co-polybuthylmetacrylate with comparable molecular weight distributions is investigated performing DSC experiments with the intention of characterize the relaxation dynamics. Results In this study PMMA homopolymer (PMMA1) and two different PMMA copolymers, namely PMMA-co-polyethylmetacrylate (PMMA2) and PMMA-co-polybuthylmetacrylate (PMMA3) having both 10 mol-% of comonomers, were used: the compositions, the molecular weights, the polydispersity index and the glass transition temperatures are reported in Table 1. Table 1. Composition and physical parameters of synthesized PMMA copolymers. sample mol-%a Mw

(Dalton) Mn

(Dalton) PI Tg (K)

PMMA1 - 59,000 35,000 1.7 398 PMMA2 10 59,000 36,000 1.6 394 PMMA3 10 70,000 47,000 1.5 385 a) Feed composition From the DSC traces obtained after a three-step thermal cycles, the peak endotherm temperature Tp is obtained as a function of annealing time as reported in Figure 1. The usual behavior of a glassy polymer can be seen here with Tp increasing linearly with log (annealing time). The only deviation from this behavior is for short annealing times for which the peak temperature appears to remain constant [1].

384

388

392

396

400

404

408

2.1 2.3 2.5 2.7 2.9 3.1 3.3

log [ta (min)]

Tp (K)

Figure 1. The dependence of Tp respect to the annealing time for PMMA1 (♦), PMMA2 (■) and PMMA3 (▲). As a general procedure in calorimetric experiments, the enthalpy lost on aging a glass for a period of time ta at a given temperature Ta can be evaluated according to the relation (1):

{ }∫ ∂−=∆y

x

T

T

up

apaa TTCTCtTH )()(),( (1)

where Cpa(T) and Cpu(T) are the heat capacity measured after the annealing and that of the unannealed sample, while Tx and Ty are the reference temperature (Tx<Tg<Ty) where the two signals overlap [2]. This procedures provides the experimentally enthalpy difference after two thermal treatments, with and without annealing. In Figure 2.a a schematic behavior of ∆H (Ta, ta) as a function of log ta is plotted and the relaxation rate and the plateau value which characterized the enthalpy relaxation process are shown. The relaxation rate RH (Ta) of the enthalpic relaxation process is defined as the inflectional slope of the relaxation isotherm [3]:

inflog),()(

a

aaaH td

tTHdTR ∆= . (2)

This parameter turns out useful in comparing the relaxation kinetics of our glass forming systems, characterized by different chain flexibility. In Figure 2.b the experimental results concerning ∆H (Ta, ta) vs log ta for three samples are shown and the results of the linear interpolation are reported in Table 2.


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∆H

log ta

∆H

log ta a)

1.8

2.2

2.6

3.0

3.4

3.8

4.2

2.1 2.3 2.5 2.7 2.9 3.1 3.3log [ta (min)]

∆H (J/g)

b) Figure 2. (a) Schematic behavior of the enthalpy lost on aging a glass as a function of annealing time; inflectional slope, defining the relaxation rate RH, and the plateau value of ∆H are shown. (b) Enthalpy lost on annealing at Ta=Tg-10°C as a function of annealing time evaluated using equation (1) for PMMA1 (♦), PMMA2 (■) and PMMA3 (▲).

As RH is a function of time, for the time range explored we are not able to say if the annealing process stays within the “intermediate” range of overall sigmoidal behavior or not; for this reason the evaluated RH (Ta) are discussed as apparent relaxation rate. As reported in Table 2, on going from PMMA1 to PMMA3, the apparent relaxation rate RH (Ta) of the relaxation isotherm at Tg-10°C, increases as expected from an increases of chain flexibility due to the presence of polyethyl and polybuthyl matacylate. Table 2.. Apparent relaxation rate RH

sample RΗ (J/g) R2

PMMA1 1.59 0.999 PMMA2 1.82 0.992 PMMA3 1.84 0.998

Conclusions In this work a study carried out by means of DSC experiments on enthalpy relaxation of PMMA copolymers has been presented. The effect of the chain flexibility arising in polymeric dynamics was highlighted evaluating the enthalpy apparent relaxation rates. References [1] J.M. Hutchinson, M. Ruddy, J. Polym. Sci. Polym. Phys.Ed. 1990, 28, 2127. [2] I.M. Hodge, J. Non-Cryst. Solids 1994, 169, 211 [3] A.J. Kovacs, J. Polym. Sci. 1958, 30, 131.


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STRUCTURE AND SUPRAMOLECULAR ORGANIZATION OF SUBSTITUTED POLYTHIOPHENES

A. Famulari, P. Arosio, G. Raos, M. Moreno, S.V. Meille

Dipartimento di Chimica, Materiali ed Ingegneria Chimica del Politecnico di Milano, via Mancinelli 7, I-20131 Milano, Italy; e-mail: [email protected]

Introduction Substituted thiophene systems are a class of stable and processable polymers whose electronic and optical properties can be tailored by monomer functionalization. The more studied are poly(3-alkylthiophenes) (P3AT’s) whose basic crystalline organization features have been established by the work of many groups.1 There is consensus on a model in which parallel stacks of polyconjugated main chains organize in layers, with the side chains running in the regions between the stacked main chains. Fitting the models to diffraction data in various instances required doubling the interlayer lattice periodicity (a in Figure 1a). However disorder, poor regioregularity and the scarce propensity to crystallization of the rigid polythiophene chain have so far prevented the actual solution and refinement of the crystal structures of these macromolecules.

Figure 1. a) Projection of the packing widely accepted for P3AT’s: in form I polymorphs chains organize in layers characterized by the a periodicity; within a layer the main chains (thick lines in projection) are stacked approximately parallel, with periodicity 1/2 b. The poly-3-(S)-2-methylbutylthiophene (b) and the poly-3-butylthiophene (c) monomer units. In the present study, we report X-ray diffraction studies2 concerning the crystal structure of poly-3-(S)-2-methylbutylthiophene (PMBT, Figure 1b) together with the investigation of preferred orientation of the poly-3-butylthiophene (PBT, Figure 1c) in solution cast films. These systems are models for crystal structure determination of substituted polythiophenes with longer side chains. Fundamental structural studies are of key importance in this field. When the particular case of enantiopure side chains are considered, chiral polythiophenes such as PMBT are obtained, which are of interest because of their potential as nonlinear optical, circularly polarized luminescent and enantiosensor materials. Meijer et al.3 found that chiral substituents can induce high optical activity in the π−π* transition of the backbone, when polymers are in aggregated, plausibly well ordered, states. However, the supramolecular and molecular organization of these aggregates were discussed on the basis of optical features without detailed characterisation.

Materials and sample preparation Powder X ray diffraction data were obtained with samples pressed in glass capillaries. Freestanding solution cast films (0.1mm thick) were prepared by slow evaporation of the solvent (chloroform). WAXD patterns were collected using graphite monochromated Cu-aK radiation, (l=1.54179Å), on a Bruker P4 diffractometer equipped with a HiStar 2D detector (sample-detector distance of about 10 cm). SAXS patterns were obtained on a Bruker Nanostar system, equipped with mirror focusing and a HiStar 2D detector (sample-detector distance of about 60cm). Results and discussion The investigation of PMBT structure have taken full advantage of the relatively simple phase behaviour (the absence of mesophase) and of the chirality of the system. The latter feature limits possible chain symmetry and simplifies the space groups analysis. The powder X-ray diffraction data (Figure 2a) allowed to propose an orthorhombic lattice with a = 13.3Å, b = 7.70Å, c (chain axis) = 7.76Å. With four monomer units, i.e. two chains, in the unit cell a calculated density of 1.27 Mg/m3 is obtained. These data are consistent with an approximate structural model of the kind suggested for form I polymorphs of P3AT’s1,2 with a somewhat expanded stacking periodicity.

Figure 2. X-ray diffraction pattern of crystalline PMBT: (a) observed profile; (b) calculated profile in C2221; (c) difference curve. The chiral side-chain, under the equivalence postulate of chemical repeating units, limited the symmetry of transoid crystalline chains to twofold screw conformations. Molecular models were obtained imposing the 21 symmetry with a 7.76 Å periodicity and exploring the potential energy surface of the infinite

a

b

a b c

a b c

10 20 30 40 2θ° 50


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a

c

isolated main chain by molecular mechanics calculations. The C2 and C2221 space groups produced the most favourable crystal packing and the corresponding structure were refined by the Rietveld technique. An high agreement between calculated and observed diffraction patterns was obtained for the C2221 space group with lattice parameters refined to the values: a = 8.24(2) Å, b = 7.75 (1) Å, c = 27.00(2)Å yielding a calculated density of 1.18 g/cc. In Figure 2 the observed (a) and calculated (b) profiles are reported together with the difference curve (c). Views of the refined molecular conformation and of the packing are shown in Figure 3.

Figure 3. Crystal conformation of PMBT refined in C2221 and molecular packing viewed along the b axis (only non-hydrogen atoms are shown) Layers arise stacking iso-directional chains: non-interdigitation of the side chains of different layers is apparent in the refined structure confirming a feature assumed in form I polymorphs of polythiophenes (Figure 1a). The small deviations from planarity of the polyconjugated main chain are hardly significant, however the refined molecular model is clearly chiral since the ordered, optically active side chains create a chiral envelope around the main chain. 2D WAXD of freestanding PBT films (see Figure 4a) obtained by slow casting from chloroform solutions show the co-presence of both form I and form II polymorphs. The signals regarding the crystalline form I (Figure 1a) present a remarkable anisotropy. The interlayer a axis is perpendicular to the film surface: as in the case of longer side chain systems, first and second order (h00) reflections are clearly confined in the equatorial direction of the 2D pattern. The scan along this direction allows the determination of the full width at half maximum (FWHM) of these peaks with a coherence length estimated to be in the range 85-90Å. 2D SAXS patterns (see Figure 4b) resulted to be anysotropic. Data analysis evidence a meridional electron density modulation comparable with the WAXD calculated coherence length.

Figure 4. 2D transmission WAXD (a) and SAXS (b) patterns of 0.1µm thick solution cast PBT film. The use of a semi-quantitative approach (Is2 vs s plot) assuming a lamellar model of alternating crystalline and amorphous regions (see Figure 5) showed peaks referred to an electron density periodicity L of about 95Å consistently with a crystalline phase long 6-8 times the interlayer main chains periodicity and in agreement with WAXD results.

Figure 5. Equatorial (orthogonal to the PBT film surface) scansion of SAXS pattern shown in Figure 2b. Conclusions In the present study we reported the solution and detailed refinement of the crystal structure of PMBT. This represents, to our knowledge, the first instance in the case of a P3AT. The simplicity of the refined structural models may result of key importance in the exploration by molecular dynamics of aggregation processes of polythiophenes as well as of disorder features and surfaces in polythiophenes aggregates. Significant anisotropy is observed by 2D WAXD and SAXS patterns for the form I of PBT in solution cast films. References 1. a) M. J. Winokur, P. Wamsely, J. Moulton, P. Smith, A. J. Heeger Macromolecules 24, 3812 (1991); b) S.V. Meille, V. Romita, T. Caronna, A.J. Lovinger,M. Catellani, L. Belobrzeckaja,. Macromolecules 30, 7898 (1997). 2. a) M. Catellani, S. Luzzati, F. Bertini, A. Bolognesi, F. Lebon, G. Longhi, S. Abbate, A. Famulari, S. V. Meille, Chem. Mat. 2002, 14, 4819; b)P. Arosio, A. Famulari, M. Catellani, S. Luzzati, S.V. Meille, Macromolecules 40 , 3-5 (2007). 3. B. M. W. Langeveld-Voss, R. A. J. Janssen; E.W. Meijer, J. Mol. Struct. 521, 285 (2000).

XX rraayy aa

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APPLYING FAST HEATING RATES

A. Fortunato1, M. Wagner2, J.E.K Schawe2 1 METTLER TOLEDO S.p.A., Via Vialba 42, 20026 Novate Milanese (MI), E-mail: [email protected]

2 METTLER TOLEDO GmbH, Sonnenbergstrasse 74, CH-8603 Schwerzenbach, Switzerland Introduction Measurements at high heating rates lead to • much shorter experiment times • increased sample throughput • better sensitivity with certain effects (e.g. the

determination of weak glass transitions) because of the greater heat flow signals at higher heating rates

• less reorganization in the sample during the DSC measurement. In order to achieve correct results a DSC should work according to the theoretical expected behaviour when working with low and high heating rates which is: Glass Transition: • ∆Cp is constant • Tg increases with heating rates Recrystallization: • Temperature increases with heating rate • ∆Hc decreases with heating rates Melting: • Peak temperature follows Illers’ law The instrumentation used for this application is DSC823e with FRS5 MultiSTARe sensor, Intracooler and Sample Changer. Crucibles used: 20 µl aluminum light pans. Due to the integrated FlexCal® calibration system, onset temperatures and the enthalpies of fusion are independent of the heating rates. This ensures that the DSC measures reliably at high heating rates and that accurate temperature and heat flow values are obtained. In summary, one can say that: • onset temperatures can be determined accurately • the same calibration procedures as with conventional

DSC (heating rate to 20 K/min) can be used.

Results Amorphous PET exhibits cold crystallization and the heating rate determines how much crystalline material is formed and hence the degree of melting that takes place. In Fig.1 the curves measured at lower heating rates do not characterize the original material because the main effects (crystallization and melting) originate during the actual measurement itself. The curve measured at 200 K/min shows only weakly induced processes and a very small melting peak. The absence of the melting peak at 240°C on the curve measured at 300 K/min shows that in this case no cold crystallization at all has occurred. This curve characterizes the original material with its original crystallinity. ∆Cp and glass transition temperatures at various heating rates follow the expected theoretical behaviour, while melting peaks of semicrystalline PET follow Illers law: Tm is proportional to (m β)1/2 .

Fig. 1: DSC curves of amorphous PET at different heating rates (specific heat capacity). An other example about liquid crystal transitions in azoxydianisole is shown. Conclusions DSC823e can measure samples at very high heating rates. FlexCal® technology ensures that the measured temperatures remain independent of the heating rate (no special calibration required). High heating rates permit undesired processes occurring during the measurement to be suppressed, so that the measurement curve characterizes the original material with shorter scan times and smaller sample mass. Advantages are: • Suppression of kinetically controlled effects (e.g. cold crystallization measurement of the

original crystallinity). • Simulation of production conditions. • Signal amplification due to the more rapid heating

rate for small samples and/or samples that exhibit only weak transitions.

References [1] Wagner M., Schawe J.E.K., Bottom R., Larbanois P., USERCOM 19, pag. 1-5, (2004) [2] Schawe J.E.K., Journal of Applied Polymer Sciences (submitted), (2004)


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INFLUENZA DELLA MICROSTRUTTURA SUL COMPORTAMENTO TERMICO DI COPOLIMERI ETILENE/4-METIL-1-PENTENE

S. Losio1, A. Abete1, S. Scafati Taglialatela1, P. Stagnaro2, G. Costa2, M.C. Sacchi1

1 ISMAC-CNR, Via Bassini 15, 20133 Milano e-mail: [email protected] 2 ISMAC-CNR, Sez. Genova, Via De Marini 6, 16149 Genova

Introduzione I copolimeri etilene/4-metil-1-pentene (E/4M1P) sono considerati materiali estremamente interessanti per il “food packaging” sin dagli anni settanta. Un limite alla loro applicabilità era dovuto al fatto che, essendo prodotti con i catalizzatori Ziegler-Natta tradizionali, erano caratterizzati da un basso contenuto di comonomero e da una inevitabile eterogeneità sia nella composizione del copolimero che nella distribuzione dei pesi molecolari. Lo sviluppo dei catalizzatori metallocenici, negli anni ottanta, e dei cosiddetti complessi “constrained geometry” (CGCs), nel corso degli anni novanta, ha dato origine alla possibilità: i) di produrre copolimeri a distribuzione omogenea dei comonomeri e ii) di controllare il contenuto di comonomero anche con olefine ramificate ad alto ingombro sterico, quale il 4M1P. Queste novità rendono possibile modulare le proprietà microstrutturali e molecolari, giocando sul contenuto e la distribuzione delle ramificazioni nella catena e sui pesi molecolari. Per questa via possono essere modulati di conseguenza anche il volume libero, il comportamento termico e, più in generale, le proprietà finali dei materiale. Lo scopo del presente lavoro è di investigare l’effetto della struttura del catalizzatore sulle caratteristiche microstrutturali di copolimeri E/4M1P e la ricaduta di queste sul comportamento termico del materiale finale. In particolare sono qui presentati i risultati del frazionamento termico a partire dal fuso condotto mediante tecnica SSA (Successive Self-nucleation and Annealing).1 Risultati A tutt'oggi non sono molti gli studi microstrutturali di copolimeri a base etilenica con comonomeri ramificati.2,3

In questo lavoro, riportiamo un’ampia analisi 13C NMR di tre famiglie di copolimeri E/4M1P ottenute utilizzando tre sistemi catalitici a diversa regio- e stereospecificità (Figura 1):2 r-CH2(3-t-BuInd)2ZrCl2 (I), r-Et(H4Ind)2ZrCl2, (II), e Me2Si(Me4Cp)(t-Bu-N)TiCl2 (III).

Figura 1. Sistemi catalitici impiegati

In precedenza il catalizzatore I, il più ingombrato ed il più stereo- e regiospecifico dei due catalizzatori metallocenici, si era dimostrato estremamente utile nell’assegnazione microstrutturale di questa famiglia di copolimeri.3 Il catalizzatore III, del tipo cosiddetto "constrained geometry", aspecifico e regioirregolare, è di grande interesse industriale per la peculiarità di abbinare all’elevata attività l'alto livello di incorporazione di α-olefine superiori. L’assegnazione dettagliata e completa che abbiamo proposto offre la possibilità di determinare in maniera corretta il contenuto e la distribuzione dei comonomeri. E’ stata valutata la distribuzione delle triadi per le tre serie di copolimeri sintetizzati e i dati microstrutturali sono stati analizzati utilizzando un metodo che permette di determinare i rapporti di reattività, r1 e r2, e il loro prodotto, r1r2, per ciascun catalizzatore.4

Tabella 1. Rapporti di reattività relativi ai tre catalizzatori

Catalizzatore r1 ± δr1 r2 ± δr2 r1r2 ± δ(r1r2)a

I 92.1±15.4 0.18±0.06 16.5± 7.9 II 100.6±14.5 0.10±0.04 10.0± 5.5 III 7.5 ±1.3 0.11 ± 0.04 0.81 ± 0.43

a δ (r1r2) = r1δr2 + r2δr1 I dati in Tabella 1 mostrano che, con la coppia di monomeri E/4M1P, i tre catalizzatori usati presentano statistiche di copolimerizzazione molto diverse. Il catalizzatore III produce copolimeri sostanzialmente random (r1r2 ≈ 1), mentre i due catalizzatori metallocenici (I e II), producono copolimeri tendenzialmente bloccosi (r1r2 » 1). Ciò che è singolare è che la bloccosità deriva non solo dall’elevato valore di r1; anche r2 ha un valore relativamente elevato, paragonato ai valori comunemente osservati con i catalizzatori “single site” nella copolimerizzazione etilene-propilene in condizioni analoghe.5 Per analizzare l’effetto delle differenze microstrutturali sul comportamento termico dei materiali si è condotto un frazionamento termico mediante SSA su alcuni campioni significativi. Questa tecnica consiste nel sottoporre il campione ad una serie di cicli di riscaldamento, raffreddamento e successivo riscaldamento ripetuti in sequenza raggiungendo ad ogni ciclo temperature massime via via decrescenti. Il tutto viene eseguito in un normale calorimetro a scansione differenziale (DSC), utilizzando quindi pochi milligrammi di materiale, con indubbi vantaggi, anche in termini di tempo e semplicità, rispetto ai tradizionali metodi di frazionamento in soluzione.

Stereo- e regiospecificità

ZrCl2 ZrCl2 SiN

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Cl

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Al termine della procedura si ottiene una profilo finale di fusione del tipo di quelli mostrati in Figura 2. Il polimero frazionato mostra una serie di picchi corrispondenti alla fusione di cristalliti costituiti, nel caso del polietilene, da sequenze metileniche di diversa lunghezza. Figura 2. Curve di fusione dopo frazionamento termico mediante SSA di due copolimeri a stesso contenuto di 4M1P (ca. 4 mol%) ottenuti dai catalizzatori II e III. Come si può osservare dai profili riportati in Figura 2, i due copolimeri contenenti ca. 4 mol% di 4M1P ottenuti con i catalizzatori II e III, pur avendo circa la stessa cristallinità complessiva, presentano distribuzioni estremamente diverse delle lunghezze delle sequenze cristallizzabili. In particolare, il catalizzatore III, con tendenza a dare inserzione casuale dei comonomeri, mostra un profilo di distribuzione delle sequenze che più si avvicina a quello di una gaussiana, con massimo centrato a 104°C, corrispondenti alla fusione di sequenze costituite da ca. 60 unità metileniche. Il catalizzatore II, che invece ha tendenza alla bloccosità, mostra

un’evidente prevalenza di sequenze che fondono a 114°C, corrispondenti a lunghezze di ca. 90 unità metileniche; in questo caso le sequenze a lunghezza maggiore sono in quantità del tutto trascurabile, mentre la quantità di quelle a lunghezza inferiore decresce più rapidamente rispetto a quanto osservato per il catalizzatore III. Conclusioni In base alla scelta del catalizzatore è possibile introdurre ampie variazioni nelle caratteristiche microstrutturali di copolimeri E/4M1P. Il frazionamento termico mediante tecnica SSA ha permesso di evidenziare le importanti ripercussioni che, a parità di contenuto di comonomero, le caratteristiche microstrutturali hanno sulla distribuzione delle lunghezze delle sequenze metileniche nei copolimeri. Ringraziamenti Questo lavoro ha avuto il supporto della Regione Lombardia. Progetto n° 308727 – INOPACK, 2005 – 2006. Gli autori ringraziano il Dr. L. Resconi per aver fornito il catalizzatore r-CH2(3-t-BuInd)2ZrCl2 e il Sig. V. Trefiletti per gli esperimenti DSC. Riferimenti 1) Müller, A.J.; Arnal, M.L. Prog. Polym. Sci., 2005, 30, 559 2) Kimura, K.; Yuasa, S.; Maru, Y. Polymer, 1984, 25, 441 3) Losio, S.; Tritto. I.; Zannoni, G.; Sacchi, M.C. Macromolecules, 2006, 39, 8920 4) Galimberti, M.; Piemontesi, F.; Fusco, O.; Camurati, I.; Destro, M. Macromolecules, 1998, 31, 3409 5) vedi ad es.: Galimberti, M.; Piemontesi, F.; Mascellari, N.; Camurati, I.; Fusco, O.; Destro, M. Macromolecules, 1999, 32, 7968

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THERMAL FRACTIONATION OF PROPENE/1-PENTENE RANDOM COPOLYMERS BY SUCCESSIVE SELF-NUCLEATION AND ANNEALING

F. Azzurri1, G.C. Alfonso2, L. Boragno1,2a, V. Trefiletti2a, F. Forlini2b, P. Stagnaro2a

1 Università degli Studi di Genova, Dipartimento di Chimica e Chimica Industriale, Via Dodecaneso 31, 16142, Genova; e-mail: [email protected]

2 Consiglio Nazionale delle Ricerche, Istituto per lo Studio delle Macromolecole, ISMAC-CNR: a Via de Marini 6, 16149 Genova; b Via Bassini 15, 20133 Milano

Introduction The crystallization and melting behavior of random copolymers of propylene and higher 1-olefins are greatly affected by nature and content of the co-units, as well as by details in their microstructure related to the used the catalytic system.1 Concentration and type of comonomer, in combination with thermal treatments, dictate the resultant structure: α-modification is largely predominant in the samples crystallized upon fast cooling from the melt. However, when crystallization is performed in isothermal conditions at relatively low undercooling, considerable amounts of γ-form can develop. Quite interestingly, for a sample containing 1-butene co-units, we found a novel reflection at 10.2° 2θ in the WAXD pattern, which has been ascribed to a new modification of i-PP.1 Actually, the formation of a new crystalline modification, which is isomorphous to form-I of it-poly(1-butene), was recently found for i-PP containing 10-25 mol% of 1-hexene.2,3 We have now clear evidence that also propene/1-pentene copolymers crystallize in this new trigonal modification over a wide range of compositions.4 Here we present the results of thermal fractionation experiments on a series of random propene/1-pentene copolymers at increasing 1-pentene content achieved applying SSA (Successive Self-nucleation and Annealing) technique.5

Thermal fractionation techniques based on the use of DSC (Differential Scanning Calorimetry), such as SC (Step Crystallization) and the above mentioned SSA are powerful methods to fractionate polymers from the melt according to the crystallizability of chain segments. They offer a simple and quick alternative to more complicated and time consuming techniques, that require to perform preparative or analytical fractionation in solution. Experimental SSA technique consists of successive self-nucleation and annealing steps carried out with a DSC instrument. Samples were kept at 220°C for 5 min, to erase any previous thermal history, then cooled (tipically at 10°C/min) up to a minimum temperature (e.g. 20°C), which is related to the starting of the melting endotherm in a “normal” DSC heating run. The cooling curve is recorded to get the standard crystallization temperature, TC. After holding the samples at the chosen minimum temperature, typically for 5 min, they were heated (at the same rate of 10°C/min) up to a selected self-nucleation temperature, called TS, located slightly above the temperature at which the melting endotherm ends. Samples were kept at TS for 5 min, then cooled again to the minimum temperature. For poorly

crystalline systems may be necessary hold the specimen at the minimum temperature for longer times, in order to achieve complete crystallization. These steps were repeated several times, each time reaching a maximum temperature which is 5°C lower than in the previous cycle, until the melting temperature range was fully covered. The “final melting” curve, which concludes the procedure, figures out the resultant thermal fractionation. An example of SSA thermal program is sketched in Figure 1. Figure 1. Temperature program for a SSA procedure. Results SC is based on isothermal crystallization steps that are repeated at progressively lower temperatures. SSA method, instead, is based on sequential application of self-nucleation and annealing steps. In the case of i-PP, regio, stereo and constitutional defects, which are present along the macromolecular chain, interrupt the crystallizable sequences. Thus fractionation takes place according to the length of regular sequences. After some preliminary experiments, the SSA method has been chosen because of its better resolution and shorter experimental time in respect to the SC method. The propylene/1-pentene copolymer samples that were fractionated by SSA are listed in Table I. The final melting curves obtained after SSA procedure clearly show a sequence of well separated peaks, each one corresponding to a set of crystallizable sequences sharing the same average length.

Table I.

Copolymer 1-pentene (mol%)

T min (°C)

T range (°C)

1 1.1 20 50-150 2 10.6 20 30-100 3 18.0 0 25-80 4 40.7 20 30-55

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Sample 1, which contains a very low amount of comonomer units, crystallizes in the classical α-form. The other samples, richer in 1-pentene co-units, also crystallize, but mainly in the new trigonal modification, which melts at lower temperature with respect to α-polymorph and slowly develops on aging at room temperature.4 As an example, the final melting curve of sample 2, containing ca. 11 mol% of 1-pentene, obtained after the SSA fractionation procedure is shown in Figure 2. The multimodal character of the standard melting curve, reported as well, is maintained; however each broad endotherm has been separated in a series of narrow peaks. Figure 2. Comparison between the final melting curve obtained after SSA fractionation and the standard melting curve for sample 2. Through a deconvolution procedure the components of the complex thermal curves have been quantitatively evaluated. One can observe from Figure 2 that SSA treatment produces a shift of the melting peaks towards higher temperatures; furthermore an increasing of the enthalpy related to the endotherm series in the temperature range 40-70°C and ascribable to the melting of the trigonal polymorph is evident. While the enthalpy value of the set of endotherms at higher temperature, which can be

ascribed to defective α- or γ-form crystals, are less affected by thermal treatment. Further investigations on SSA fractionated samples are currently in progress, also by using temperature resolved SAXS/WAXD techniques. Conclusions A series of propylene/1-pentene random copolymers at increasing content of 1-pentene obtained from a metallocene catalyst were fractionated from the melt by using successive self-nucleation and annealing technique. A deeper insight into the microstructure of the samples, in terms of length distribution of the crystallizable stereo-sequences and role of 1-pentene co-units in developing the new trigonal modification of i-PP was achieved. Acknowledgements Authors warmly thank Mrs. V. Pierantoni for performing DSC experiments. References 4. P. Stagnaro, G. Costa, V. Trefiletti, M. Canetti, F. Forlini, G.C. Alfonso, Macromol. Chem. Phys., 207, 2128 (2006) 5. C. De Rosa, F. Auriemma, P. Corradini, O. Tarallo, S. Dello Iacono, E. Ciaccia, L.Resconi J. Am. Chem. Soc., 128, 80 (2006) 6. B. Lotz, J. Ruan, A. Thierry, G.C. Alfonso, A. Hiltner, E. Baer, E. Piorkowska, A. Galeski Macromolecules, 39, 5777 (2006) 7. G.C. Alfonso, F. Azzurri, L. Boragno, G. Costa, M. Canetti, P. Stagnaro contribution submitted to this conference 8. A.J. Müller, M.L. Arnal, Prog. Polym. Sci., 30, 559 (2005)

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SUPRAMOLECULAR CHIRALITY IN NEW CHIRAL LIQUID CRYSTAL AZOPOLYMERS OBTAINED BY ATRP

L. Giorgini1, L. Angiolini1, J. Barberà2, F. Paris1, E. Salatelli1, R. M. Tejedor2

1 Università degli Studi di Bologna, Dipartimento di Chimica Industriale e dei Materiali, INSTM UdR di Bologna, Viale Risorgimento 4, 40136-Bologna, Italy e-mail: [email protected]

2 Química Orgánica, Instituto de Ciencia de Materiales de Aragòn, Facultad de Ciencias-ICMA, Universidad de Zaragoza-CSIC, Pedro Cerbuna 12, 50009-Zaragoza, Spain

Introduction Side-chain liquid crystal polymers (LCP) with tailored azobenzene moieties are widely studied due to the easy photomodulation of their properties (birefringence, non-linear optical properties and chiroptical properties) making them suitable for several applications as materials for holographic data recording, non-linear optics, artificial muscles or molecular switches 1-5.The presence of a chiral group of one absolute configuration in the side chain can lead to induction of supramolecular chirality both in amorphous and liquid crystal state. Results Here, we present the synthesis by ATRP and the characterization of monodisperse chiral azopolymers of well-controlled structure (linear and three-arms star shaped) and different average molecular weights. These new azopolymers include L-lactic residues in the side chain, which can transmit their chirality to a supramolecular organization of the liquid crystal phase. The nature of the liquid crystal phases of the synthesized polymers has been characterized by DSC, powder X-ray diffraction (XRD) and polarized optical microscopy (POM).

Fig 1: Structure of the new monomer (S)-ML6A, linear and star shaped polymers. The XRD results show the presence of a SA1/2 phase with a POM texture typical of helical TGB phase. In the SA phase, where the molecules are arranged perpendicular to the layer planes, the molecules cannot adopt a helical structure perpendicular to the layers. It can only occur parallel to the layers and only for systems with a strong twisting power. Therefore, a helical superstructure is only possible if screw dislocations punctuate the layers, giving rise to the so-called twisted grain boundary A phase (TGBA).6

Photomodulation of chiroptical properties On other hand, it is well-know that is possible to photomodulate the chiroptical properties of thin films of azoaromatic polymers by irradiation with circularly polarized (CP) light of one single left (L) or right (R) rotation sense. The chiroptical properties of thin films of the investigated LC polymers were studied before and after irradiation with R- and L-CP light by circular dichroism (CD). The annealed thin films of polymers in TGB phase were irradiated with CP-R and CP-L, respectively, with an argon laser at 488 nm (I power = 25 mW/cm2) for 30 min. The UV-vis spectra of irradiated polymers result similar to the spectra of annealed corresponding films. The CD spectra of these polymers after a cycle of illumination with CP-R and CP-L light are shown in Figure 2. By irradiation with CP-R radiation, the CD spectra of the initially chiral annealed thin films of polymers in TGB phase display a net inversion of sign and a relevant amplification of absolute chirality. Irradiation with L-CP light a other native films for 30 min afforded an induced optical activity of similar magnitude to that observed after irradiation with R-CP, but with the opposite sign.

200 300 400 500 600-2500

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Fig 2: CD spectra of films of Poly[(S)-ML6A]-14 (straight line) and Star[(S)-ML6A]-24 (square) irradiated with L-CP (___) or R-CP (····). The inversion of the couplet can be clearly seen. The resulting spectra actually appear as mirror images of each other both for linear and star polymers. The CD spectra of the irradiated linear polymer display higher ellipticity values than the CD spectra of the irradiated star ones. The observed effects are reversible; when the handedness of the pump beam was switched from right to left and the irradiation was performed on the same illuminated films for 30 min, a similar opposite CD spectra were obtained (Figure 2).

O

O

Br

O O

OO

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NN

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Monomer Linear Polymer Star Polymer

OO

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In conclusion, the experiments suggest that R-CP light induces right-handed supramolecular conformation of the material and the resulting helix has been correspondingly erased and reinscribed with left CP radiation. The induced circular anisotropy can be erased by heating the irradiated samples on their correspondents isotropization temperatures. Discussions The mechanism of the reversible helical inversion induced by CP radiation is not well understood. In any case, for the investigated polymers, it appears related to a preliminary helix supramolecular ordering of the azobenzene molecules. In fact, no reproducible circular anisotropies can be photoinduced in the native films (not annealed), pointing out the essential role of the liquid crystalline arrangement. This also indicates that orientational preorganization is necessary for obtained a controllable photomodulation of chirality. It is known that CP light tends to align the azobenzene side groups along directions close to normal to the film.5 It is possible that transfer of angular momentum from the CP radiation to the medium, as occurs when a CP photon is absorbed, induces a precession of the chromophores with a sense of rotation congruent with the sense of the CP light. This would mean that CP-L light induces a left-handed organisation of the azobenzene molecules, whereas CP-R light induces a right-handed one. In this mode, the sign of the CD signals associated with the conformational aggregation of neighbouring chromophores inverted as we have recently demonstrated by studies on the dimeric derivative 2,4-dimethylglutaric acid bis (S)-3-[1-(4′-nitro-4-azobenzene)pyrrolidine ester, corresponding to the smallest section of the polymeric chain where interchromophore interactions are relevant6,7. Its CD spectrum shows an exciton couplet of strong amplitude which suggests that the chiral interactions between a couple of chromophores in solution are already important and that the optical activity of these materials should be substantially related to relatively short chain sections with conformational dissymmetry of one prevailing screw sense. In the case of investigated polymers, it is reasonable to think that the TGB phase, in the native annealing films, is composed of several layers disposed in a helical conformational supramolecular structure, where the polymer backbones exist mainly in the planes defining the smectic A layers, but can also go between the various smectic grains, with a prevailing twist sense due to the thermodynamically favoured chiral interaction between neighbouring l-lactic-azoaromatic moieties, thus conferring a prevailing absolute chirality to the materials. The helix direction is along the film thickness (as well as the light propagation direction), and the azoaromatic groups are mainly parallel to the film plane.

It seems reasonable that the CP light should be able to alter the chiral interaction between the optically active chromophores and consequently the helical supramolecular structure so as to reverse the absolute macroscopic chirality of the materials. For the present polymers, the interaction with CP-R light does not affect the signs of the CD signals (we observe a amplification of chirality of the same sign), whereas CP-L light leads to their net inversion, thus allowing the assignment of a prevailing right-handed conformation to the native liquid-crystal TGB phase. This supramolecular change not perturbs the texture (TGB) phase of the investigated polymeric films as shown by POM analysis of the irradiated area, before and after application of CP light.

Conclusions All the aforementioned CD effects persist for at least one month at room temperature and are well reproducible. Clearly, for technological applications, also the switching time would be important. Further investigations of the dynamics of the observed phenomena are in progress. These experiments demonstrate that it is possible to control and tune the supramolecular chirality of a polymeric liquid crystal TGB phase using a appropriate handedness of the incident CP radiation. To our best knowledge, these are the first examples of reversible chiroptical switching between two enantiomeric suprastructures obtained by a native chiral liquid-crystal TGB phase. The switching of the chiroptical proprieties made these materials suitable to be used as chiroptical switches and of potential interest in nanoscale technologies for all-optical manipulation of informations. References 1 T. Ikeda, O. Tsutsumi, Science 268, 1873 (1995) 2 S. Xie, A. Natansohn, P. Rochon, Chem. Mater., 5, 403 (1995). 3 T. Verbiest, M. Kauranen, A. Persoons, J. Mater. Chem. 9, 2005 (1999). 4 R. M. Tejedor, M. Millaruelo, L. Oriol, J. L. Serrano, R. Alcalà, F. J. Rodrìguez, B. Villacampa, J. Mater. Chem. 16, 1674 (2006). 5 G. Iftime, A. Lagugnè, F. Labartet, A. Natansohn, P. Rochon, J. Am. Chem. Soc., 121, 12646 (2002). 6 J. W. Goodby, in Structure and Bonding, Vol. 95, Liquid Crystals II, ed. D. M. P. Mingos, Springer, Berlin, 1999, p. 83. 7 L. Angiolini, T. Benelli, L. Giorgini, A. Painelli, F. Terenziani, Chem. Eur. J. 2005, 11, 6053-6353.


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NUOVO MECCANISMO DI DEGRADAZIONE TERMICA DI POLI(3-IDROSSIBUTIRRATO) INDOTTA DA GRUPPI CARBOSSILATI

M. L. Focarete1, M. Scandola1, M. Kawalec,2 G. Adamus,2 P. Kurcok2, M. Kowalczuk2

1 Università di Bologna, Dipartimento di Chimica “G.Ciamician”, via Selmi 2, 40126 Bologna, Italy; e-mail: [email protected]

2 Institute of Polymer Chemistry, Polish Academy of Sciences , 34 M. curie-Sklodowska Street, 41-819 Zabrze (Poland)

Introduzione Il poli[(R)-3-idrossibutirrato] di origine naturale (nPHB) è un poliestere batterico ad alto peso molecolare che viene sintetizzato dai microorganismi come riserva di carbonio ed energia. Gli analoghi sintetici del nPHB possono essere sintetizzati mediante polimerizzazione per apertura d’anello del β-butirrolattone ottenendo PHB con diversa tatticità (isotattico, atattico, sindiotattico). Spezzoni di PHB a catena corta, indipendentemente dalla loro tatticità, possono essere preparati dal PHB ad alto peso molecolare mediante degradazione termica ad alta temperatura. E’ noto che il PHB degrada termicamente con un meccanismo di cis-eliminazione intramolecolare, che porta ad una scissione random della catena polimerica, con la formazione di oligomeri contenenti come prodotti principali gruppi terminali trans-crotonati.1 Questo meccanismo è stato considerato fino ad ora l’unico meccanismo di degradazione termica del PHB. In questo lavoro si riporta un nuovo meccanismo di degradazione termica del PHB, che agisce in modo sia intra- che inter-molecolare ed è attivo a temperatura relativamente bassa, in presenza di gruppi terminali carbossilato. Questo risultato è oggetto di domanda di brevetto.2

Risultati Sono stati utilizzati PHB con diversa microstruttura (isotattici e atattici), diverso peso molecolare e diverso gruppo carbossile terminale. In particolare si sono utilizzati PHB contenenti come terminale di catena un gruppo carbossilico (-COOH) o un gruppo carbossilato (-COO-) con diverso tipo di controione. L’analisi termogravimetrica (TGA) ha rivelato che la stabilità termica di PHB con gruppi terminali carbossilici non dipende dal peso molecolare e dalla microstruttura, in accordo con precedenti dati di letteratura.3 Al contrario, in presenza dei gruppi terminali carbossilato si è trovata una grande influenza sulla stabilità termica di PHB. In particolare, il PHB mostra una diminuzione significativa della temperatura di massima velocità di perdita di peso (Tmax), che cala da T = 291°C a temperature che a loro volta dipendono dal tipo di controione. A fronte dell’effetto della presenza dello ione carbossilato sulla Tmax si è osservato anche un effetto sul peso molecolare del polimero durante trattamenti termici a temperatura assai modesta. Infatti, sono stati condotti esperimenti di degradazione termica in condizioni isoterme a 120°C su aPHB contenente gruppi terminali carbossilati e tetrabutil ammonio come controione, che sono stati

confrontati con risultati di analoghe prove su aPHB con terminale COOH. Dopo 4h a 120°C si è osservata una significativa diminuzione di peso molecolare nel aPHB con terminali carbossilato, mentre l’aPHB con terminali carbossilici non ha mostrato nessun calo di peso molecolare. Questo fenomeno può essere spiegato sulla base di un nuovo meccanismo di degradazione termica di tipo E1cB, illustrato nello Schema I.4 Secondo tale meccanismo la degradazione avviene mediante una α-deprotonazione ad opera dei gruppi carbossilati e può procedere sia a livello intermolecolare che intramolecolare. Il carattere intermolecolare della reazione di degradazione termica di PHB indotta da carbossilati secondo il meccanismo E1cB è stato inoltre dimostrato da prove condotte su leghe polimeriche di PHB isotattico e PHB atattico con terminale carbossilato, che è stato dimostrato indurre degradazione termica nel polimero di origine naturale con terminale COOH.. Conclusioni In questo lavoro è stato proposto un nuovo meccanismo di tipo E1cB per la degradazione termica di PHB indotta da gruppi carbossilati. Questo è il principale tipo di decomposizione che opera a temperature moderate in questo poliestere che contiene un H in α al carbonile. I risultati mostrano inoltre che è possibile ottenere la degradazione controllata di polimeri contenenti un H in α al carbonile, ricavandone oligomeri da utilizzare come building blocks in sintesi chimiche. Riferimenti 1. N. Grassie, E.J. Murray, P.A. Holmes, Polym. Degrad. Stab. 6, 127-134 (1984) 2. M. Scandola. M.L. Focarete, I. Foltran, M. Kowalczuk, P. Kurcok, M. Kawalec, G. Adamus “Method of controlling thermal degradation of anionically terminated polymers and materials obtained thereof”, PCT Patent application (filed on March 20, 2006) at No. PCT/IB2006/000898 3. P. Kurcok, M. Kowalczuk, G. Adamus, Z. Jedliński, R.W. Lenz, J. M. S. - Pure Appl. Chem. A32, 875-880 (1995) 4. M. Kawalec, G. Adamus, P. Kurcok, M. Kowalczuk, I. Foltran, M.L. Focarete, M. Scandola, Biomacromolecules 8, 1053-1058 (2007)


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Schema I: Meccanismo E1cB di degradazione termica di PHB


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THERMAL DEGRADATION OF ANHYDRIDE CURED EPOXY RESINS

M.P. Luda1 , A.I. Balabanovich 2, D. Guaratto1 1 Università di Torino, Dipartimento di Chimica IFM, Via P. Giuria 7, 10125 Torino, Italy; e-mail

[email protected] 2 Institute for Phys. Chem. Problems of Belorussian State University, Leningradskaya 14, 220050 Minsk, Belarus.

Introduction The fire risk is a major drawback of epoxy resins which have to meet flame retardancy in most of their application. This can be accomplished by incorporating brominated reactive compounds, like DGETBBA used as co monomer with DGEBA to obtain fire retardant materials. However bromine containing structures could produce corrosive and obscuring smoke and might give super toxic halogenated dibenzodioxines (PBDD) and dibenzofurans (PBDF) when burned. Polymer wastes include epoxy resins containing brominated aromatics. Attractive economical revenue is recovering of gold from dismissed printed circuit boards by a thermal treatment of them which is one of the approaches of recycling polymer wastes. In this respect, it has been reported that brominated epoxy resins are less stable at high temperatures than the non-brominated ones with relevant effect on the set up of temperature and residence time of thermal recycling. Curing of epoxies can be carried out by a number of hardeners such as amines and anhydrides. In comparison to ammines, anhydride cured epoxy in general provides longer pot lives, measured in days or months, and mixtures with lower viscosity at the processing temperatures and cure with very slight exotherm. Results Network structure Epoxies resins have been prepared on purposes by mixing phtalic anhydride with DGEBA (eq. weight 380, (ER-PHT) or with a mixture 1/1 DGEBA/DGETBBA (BER-PHT) at 50% of stoichiometry. Curing was carried out at 120°C for 3 h.

Fig. 1 IR spectra of cured ER/PHT (upper) and BER/PHT (bottom) systems In comparison to uncured systems, the IR spectra of the cured systems in figure 1 show increasing amount of –OH (3450 cm-1), C=O (1724 cm-1), 1264 cm-1 (Ar-O in DGETBBA) 1230 cm-1 (Ar-O in DGEBA) and decreasing amount of oxirane ring (3046 and 910 cm-1).

The presence of DGETBBA in BER-PHT is also accredited by the absorption at 1462 (Brominated aromatic ring) and 735 cm-1 (1,2,3, 5 substitution). Thermal stability TGA in nitrogen of these systems (Figure 2) show that, , the main stage of weight loss the is located around 383°C in BER-PHT and around 418 °C in ER-PHT. In addition a larger amount of residue at T> 450°C is left by BER-PHT. In parallel to what found in amine hardeners cured epoxies, the presence of bromine reduces the thermal stability.

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U niversal V4.1D T A Instrum ents Fig. 2 TGA in nitrogen of cured ER/PHT(●) and BER/PHT (□) systems Volatiles and high boiling products By pyrolysing at 500°C for 15 sec. ER-PHT acrolein, and 2-propen-1-ol are produced together with higher molecular weight products such as phtalic anhydride, DGEBA and monoglycidyl ether of bisphenol A. Assignation of DGEBA structure has been supported by GM-MS of original monomer. (Figure 3) At 500°C 15 sec. BER-PHT in Py-GC-MS evolves Propene, 2-propen-1-ol, Acrolein, 2- bromo-1-propene, and traces of bromoacetone and amongst higher molecular weight products phtalic anhydride, phenol and bisphenol A and brominated (1-4)bisphenol A. DGETBBA has not been found between degradation products of such system, indicating its thermal lability(Figure 3) Residues BER-PHTand ER-PHT leaves respectively at 450°C 20% end 10% of a black, fully oxidable residue. The structure of these residues is similar and shows IR typical absorption bands of the charred residues at 1600, 870, 816 and 742 cm-1.

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Fig. 3 GC-MS of DGEBA and DGETBBA monomers and Py-GC_MS od ER-PHT and BER-PHT. The same GC programs has been used.

Discussion The structure of the network is as follows:

O CH2 CHOCO

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O CCH3

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XX

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X: H or Br Where β hydrogen are available, phtalic anhydride and epoxy monomers are formed by non radical pyrolysis of the ester groups involving a six member ring transition state. Lower molecular weight products acrolein and propenol come from further radical degradation of the epoxy monomer; While DGEBA possess some volatility which bring it far from the hot oven, the less volatile DGETBBA remain in the hot area and is subjected to a more pronounced degradation.

However the presence of brominated aliphatic in the degradation products of BER-PHT suggests that radicals formed in this stage of degradation are able to extract bromine from TBBA unit leading simultaneously to an increased amount of residue.

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PHYSICAL PROPERTIES OF CROSSLINKED ALGINATE FILMS

R. Russo 1, M. Malinconico 1, G. Santagata 1 1 Institute of Chemistry and Technology of Polymers (ICTP) CNR Via Campi Flegrei, 34 80078 Pozzuoli (NA) Italy

e-mail: [email protected] Introduction Alginates are linear water soluble polysaccharides comprising (1–4)-linked units of α-D-mannuronate (M) and β-L-guluronate (G) at different proportion and different distribution in the chain (1). They are present in brown algae and can also be found in metabolic products of some bacteria. The chemical composition and the sequence of M and G residues depend on the biological source and on the state of maturation of the plant. Sodium alginate can be considered a block copolymer containing three types of blocks (GG, MM, MG). The stiffness of the three blocks decreases in the order GG>MM>MG. The physical and chemical properties of alginates strictly depend both on their composition and on the sequence of guluronic and mannuronic residues in the polymeric chain. Alginates are well known natural ionic polysaccharides used mainly as food additives, thickener, gelling agent, and in the controlled delivery of drugs. The great interest with this family of polysaccharides is strictly related to the gelling properties. Alginate solutions can form gels either by lowering the pH below the pKa value of the uronic residue, or in presence of divalent ions. The divalent ions cooperatively interact with blocks of guluronic units to form ionic bridges between different chains. The ability of forming this kind of interaction depends on the length of the G-blocks. The most popular model to account for the chain-to-chain association is the “egg box model”. In this model the G-blocks form a three-dimensional arrangement in which Ca ions are embodied in cavities like eggs in a cardboard egg-box. (2) In this study we analyze three different alginates with different molecular parameters with respect to the different ability to ionically crosslink with calcium ions. The molecular characterization of the three alginates is reported in Table 1. From the data of Table 1 we notice that the three samples are similar in terms of the molecular mass. This is particularly interesting since all the differences in some properties can be ascribed to the different concentration of guluronic residue. The films were prepared by casting from aqueous solutions. Solutions at 1% of the three polymers were prepared, filtered and kept for a few minutes under vacuum. The films were obtained pouring the solutions on polyester mold kept in plane to ensure the homogeneous thickness of the films. The films thicknesses were in the range of 40 µm. All films were allowed to form during exposure to the atmosphere for three days. Prior to testing, the films were equilibrated at 45% relative humidity by storing them in a desiccator over a saturated solution of calcium nitrate at room temperature. The crosslinked samples were prepared by soaking the films in a 2% w/v aqueous solution of calcium chloride for 30 minutes.

Table 1 Alginates molecular parameter

SAMPLE %G Mn (Da) Mw(Da) Mw / Mn

A 37 ± 1 2.3 . 105 1.2 . 106 5.2 B 45 ± 1 1.1 . 105 5.9 . 105 5.4 C 62 ± 1 2.0 . 105 5.1 . 106 5.5

The films of plain polymers and the films crosslinked with calcium ions underwent to thermal analysis (TGA, DSC, DMTA) and to test of water permeability. Before testing, the thickness of the three samples was measured before and after the immersion in the solution of calcium chloride. The samples swelled and at the same time crosslinked. The thickness of swollen samples was measured after the drying and after conditioning at room temperature and 45% relative humidity. The thickness did not return to the original values since the crosslinking points, introduced in the swollen state, stabilize the conformations of this state. Moreover, on increasing the amount of G-units the thickness increases more ranging from the 37% in the sample A to 56% in the sample C Results and discussion In Table 2 the main results of DSC, DMA and permeability analyses are reported

Table 2 Tg, Tß,, Ea and Permeability values

Sample Tg

from DSC (°C)

Tß from

DMTA (°C)

Ea (kJ/mole)

Permeability (ng·s-1·m-

1·Pa-1) A 133 15 73 0.092 B 132 32 80 0.098 C 138 48 83 0.058

A-Ca 131 -13 75 0.133 B-Ca 127 -30 59 0.142 C-Ca 119 -35 68 0.137

Sample C has the highest value of Tg due to the higher concentration of guluronic units. We notice that the Tg decreases as a result of the crosslinking. Usually, with the introduction of the crosslinking points, the Tg increases because of the reduced mobility of chain segments. In our case the crosslinking appreciably decreases the Tg, and the decrease becomes more pronounced on increasing the quantity of G-units. The increased mobility of chain segments in crosslinked


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samples can be explained in terms of increased free volume due to crosslinking in the swollen state. Since the decrease of Tg is more pronounced on increasing the content of G-units, we think that the free volume increases more on increasing the content of G. This hypothesis is also confirmed by the swelling data. The crosslinking points represent a hindrance for the packing of chains, but at the same time, the chain segments between two consecutive crosslinking points, experience an increased mobility because of the increase of the free volume. In addition, from the thermogravimetric data, we notice that crosslinked samples contain more bound water that contributes to the lowering of the Tg. From DMTA analysis, particularly from the curves of tanδ, it is possible to detect a relaxation at temperatures lower than Tg (ß transitions). The ß transition is a broad relaxation and it is generally ascribed to the motion of short segments of chain or to lateral groups. For the alginates we do not have any support from the literature but for systems like poly(vinyl alcohol) the relaxation has been observed. According to some authors, it is related to the motion of hydroxyl groups connected to water molecules, (3) while other authors consider the ß relaxation as a typical water relaxation. In order to better classify this transition the dynamic-mechanical test was performed again in multifrequency mode. It was found that on increasing the frequency, the peak maximum was shifted towards higher temperatures and this allowed us to calculate the apparent activation energy of the process. The values of the activation energies, reported in Table 2, are compatible with the motion of hydroxyl groups. It is interesting to note that the Tß increases on increasing the content of guluronic moiety. This because, on increasing the content of G units, the stiffness of chain segments increases and the molecular mobility decreases. On the contrary, in crosslinked samples we observe a decrease of the temperature at which the transition occurs and a decrease of the activation energy going from sample A-Ca to sample C-Ca. The more pronounced depression of the Tß, observed in the two samples with higher content of guluronic units, is a confirmation of our hypothesis. In fact from the swelling data we notice that samples B and C swell much more than sample A. In addition, from thermogravimetric analysis, we deduce that sample B-Ca and C-Ca contain more water than sample A-Ca. Both factors, the increase of the free volume, along with the increased amount of water, contribute to lower either the temperature or the activation energy of this transition.

The permeability data evidence that samples A and B have a similar permeability, while sample C, containing more guluronic units, has a lower value. This is reasonable because increasing the concentration of guluronic units the chain-to-chain interactions increase and act as a hindrance to the diffusing molecules making the diffusive path more tortuous. As concerning the permeability of crosslinked samples, we measured increasing values passing from A-Ca sample to C-Ca. Usually when points of crosslinking are introduced in a polymeric film the permeability reduces. There are two effects to be considered: the introduction of crosslinking points that generally lowers the transport properties and the increase of the free volume that makes the pathway of the diffusing molecules easier with the consequent increase of the transport parameters. According to the slight increase of the permeability in the crosslinked samples, we have to conclude that the increase of free volume has more effect with respect to the introduction of crosslinking points. Conclusions In this work the physical behavior of three films of alginate has been studied. The three samples had different content of guluronic moieties. The increase of the G-units promoted the chain-to-chain interaction with a consequent increase of the Tg. The crosslinking caused a considerable change of the analyzed physical parameters in an unexpected direction. In fact, both the glass transition temperature and the ß relaxation, detected with the dynamic-mechanical experiment, decreased in the crosslinked samples. This was explained in terms of increased free volume during the crosslinking process. In fact, since the introduction of the crosslinking points occurred in the swollen state, once the absorbed water was removed, the tie-points introduced between different chains, did not allow the chain segments to restore the conformations they had before the crosslinking. This hypothesis was supported by the swelling experiments. In fact the increase of the thickness as a consequence of the crosslinking process is proof of the increased free volume. The permeability data were supportive of this hypothesis as well. References 1) K. I. Draget, G. Skiak-Braek, O. Smidsrod, Int. J. Biol. Macromol., 21, 47 (1997) 2) E.R. Morris, D.A. Rees, D. Thom, J. Boyd, Carbohydr. Res., 66, 145 (1978) 3) M. Krumova, D. Lopez, R. Benavente, C. Mijangos, J.M. Perena, Polymer, 41, 9265 (2000)


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MALDI TOF/TOF AND MALDI TOF/TOF MS/MS ANALYSIS OF SYNTHETIC POLY(ESTER AMIDE) FROM SEBACIC ACID AND 4-AMINO-1-BUTANOL

P. Rizzarelli, C. Puglisi

Istituto di Chimica e Tecnologia dei Polimeri - Consiglio Nazionale delle Ricerche; V.le A. Doria 6 - 95125 Catania e-mail: [email protected]

Introduction Matrix-assisted laser desorption/ionization time-of-flight mass spectrometry (MALDI-TOF MS) has become an important tool for the analysis of complex polymer systems as it is a highly sensitive, non-averaging, technique that provides detailed structural information about the individual molecules contained in a polymer sample. Although MALDI is a soft ionization technique, product ion analysis can be performed with a MALDI-TOF mass spectrometer by means of the so-called post source decay (PSD). PSD-MALDI has been used successfully in the characterization of biopolymers, but has been used less extensively for synthetic polymers, reasonably on account of the enhanced complexity of the samples involved. Successful attempts to simplify and improve experimental procedures, acquiring more reliable information, have been done with the development of newer instruments, such as the MALDI-Q/TOF. Another recent mass spectrometer design that overcomes many of the limitations of PSD and CID in a conventional MALDI instrument is the MALDI-TOF/TOF, in which high velocity ions produced in a high-vacuum pulsed MALDI-TOF are selected with a timed ion gate, collided at kiloelectronvolt energies with gas atoms or molecules, and then further accelerated for analysis in a two-stage reflectron TOF. Recently, Matrix-Assisted Laser Desorption/Ionization Time-Of-Flight/Time-Of-Flight Tandem Mass Spectrometry (MALDI-TOF/TOF-MS/MS) has been applied in the analysis of biopolymers, such as peptides and proteins, N-glycans, oligosaccharides, showing noteworthy and improved results with respect to those achieved by PSD-MALDI. Very recently MALDI-TOF/TOF-MS/MS has proven to be a convenient and appropriate technique to obtain structural information about the ester and amide bonds sequences in synthetic polymers.1 Furthermore, more recently MALDI-TOF/TOF-MS/MS has been dedicated to the investigation of the fragmentation pathways of poly(butylene adipate) oligomers.2 In the present study, MALDI-TOF/TOF-MS/MS was employed, as a continuation of our work, to analyze a poly(ester amide) sample (PEA-Bu) from the melt condensation of sebacic acid and 4-amino-1-butanol. In particular we investigated the fragmentation pathways, the ester/amide bond sequences in the polymer chains and the structure of species not predictable according to the method of synthesis. Results An Applied Biosystems 4800 MALDI TOF/TOFTM Analyzer (Applied Biosystems, Framingham, MA, USA) mass spectrometer was used in this study to acquire MALDI and MS/MS spectra using air or argon as the

collision gas. This TOF/TOF instrument is equipped with a Nd:YAG laser with 355-nm wavelength of <500ps pulse and 200 Hz repetition rate in both MS and MS/MS modes. The QuanTIS™ precursor ion selector of the 4800 MALDI TOF/TOFTM Analyzer has a mass resolution of about 400. All measurements were performed in automatic mode. For MS/MS experiments, the collision energy, defined by the potential difference between the source acceleration voltage and the collision cell, was set at 2 kV. MALDI-TOF/TOF-MS and MALDI-TOF/TOF-MS/MS spectra were recorded in reflector positive ion mode. MS/MS experiments were all performed under “metastable suppressor on” conditions. MS and MS/MS data were processed using Data Explorer 4.4 (Applied Biosystems). Figure 1 displays the MALDI-TOF/TOF mass spectrum in reflector mode of the synthesized PEA-Bu sample together with an enlarged portion in the m/z range 980-1160. The spectrum extends up to m/z 6000 and shows a distribution of ions corresponding to singly charged sodium and potassium adducts. The repeat unit of PEA-Bu is reported in Figure 1 and its mass (255.18 Da) matches the difference between the m/z values of species belonging to two consecutive clusters.

Figure 1. MALDI/TOF-TOF spectrum of the PEA-Bu

Thirteen MALDI peaks in the inset of Figure 1 can be assigned (Table 1) unambiguously and they are due to only four oligomers: i) sodiated and potassiated ions of linear di-carboxyl terminated chains (species A Na+, m/z 990.68; A K+, m/z 1006.66) and the corresponding sodium salt of the same species (A’ Na+, m/z 1012.67); ii) protonated, sodiated and potassiated ions of cyclic PEA-Bu chains (species B H+, m/z 1021.76; B Na+, m/z 1043.75; B K+, m/z 1059.73); iii) protonated, sodiated and potassiated ions of linear oligomers terminated with carboxyl at one end and amino alcohol groups at the other end (species C H+, m/z 1039.77; C Na+, m/z 1061.76; C K+, m/z 1077.73), and the corresponding sodium salt of the same species (C’ Na+, m/z 1198.77);

600 1680 2760 3840 4920 6000m/z0

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2082.55

—[OCH2CH2CH2CH2NHCOCH2CH2CH2CH2CH2CH2CH2CH2CO]n—

980 1016 1052 1088 1124 1160m/z

C Na+1061.76

A Na+990.68

B Na+1043.75

D Na+1132.84

X1114.83

C K+1077.73

C’ Na+1083.75

A K+1006.66

A’ Na+1012.67

B K+1059.73

B H+1021.76

D K +1148.811130.80

D H +1110.85

600 1680 2760 3840 4920 6000600 1680 2760 3840 4920 6000m/z0

10

20

30

40

50

60

70

80

90

10

20

30

40

50

60

70

80

90

100

% In

tensit

y

1061.76

1316.96

990.681572.151245.88

1827.35

2082.55

—[OCH2CH2CH2CH2NHCOCH2CH2CH2CH2CH2CH2CH2CH2CO]n—

980 1016 1052 1088 1124 1160980 1016 1052 1088 1124 1160m/z

C Na+1061.76

A Na+990.68

B Na+1043.75

D Na+1132.84

X1114.83

C K+1077.73

C’ Na+1083.75

A K+1006.66

A’ Na+1012.67

B K+1059.73

B H+1021.76

D K +1148.811130.80

D H +1110.85


- - 136

iv) protonated, sodiated and potassiated ions of linear oligomers terminated with amino alcohol groups at both chain ends (species D H+, m/z 1110.85; D Na+, m/z 1132.84; D K+, m/z 1148.81).

Table 1. Symbol Structure M•Na+ A HOCO(CH2)8CO-[U]3-OH 990.68 A’ HOCO(CH2)8CO-[U]3-O-Na+ 1012.67 B [U]4

1043.75 C H-[U]4-OH 1061.76 C’ H-[U]4-O-Na+ 1083.75 D H-[U]4-O(CH2)4NH2 1132.84

U = -O(CH2)4NHCO(CH2)8CO- or -NH (CH2)4OCO (CH2)8CO-

Noticeably, the chains terminated with amino alcohol groups can bear a hydroxyl (-NH(CH2)4OH) or an amine (-O(CH2)4NH2) end group that can not be discriminate by mass spectrometry. Oligomers of uncertain structures are observed as well (m/z 1089.16, 1096.82, and 1114.83) and their occurrence could be related to side reactions taking place in the polymerization process. In fact, 4-amino-1-butanol end groups can undergo dehydration leading to the formation of –NH(CH2)2CH=CH2 groups. MALDI-TOF/TOF-MS/MS analysis using air or argon as the collision gas was performed on sodiated oligomers terminated by di-carboxyl groups (m/z 990.68), on sodiated ions of cyclic PEA-Bu chains (m/z 1043.75), on sodiated oligomers with carboxyl and amino alcohol chain ends (m/z 1061.76) and on species of doubtful structure (m/z 1114.83).

Figure 2. MALDI-TOF/TOF-MS/MS spectrum of the sodiated dicarboxyl terminated oligomers at m/z 990.7 Figure 2 shows the MALDI-TOF/TOF-MS/MS mass spectrum, acquired using air as the collision gas, of the sodiated dicarboxyl terminated oligomers at m/z 990.7 of the PEA-Bu sample. According to the structures of the most abundant ions, the main cleavages (Scheme 1) proceed through a β-H-transfer rearrangement, leading to the selective scission of the −O−CH2− bonds. Product ions originated from the −CH2−CH2− (β−γ) bond cleavage in the diacid moiety are detected too. Their formation should be supported by the presence of a α,β−unsaturated ester or amide end group. The spectrum also displays clusters of ions at an interval of 14 (CH2) mass units that resemble the series typical of hydrocarbon and long-chain acids.

Scheme 1. Most frequently observed bond cleavages and end group structures of product ions in PEA-Bu

The structure of the repeat unit presented in the figure 1 is not necessarily representative of the actual bond sequences in the poly(ester amide) chains. It was arbitrarily chosen to report the ester/amide sequences as alternate, but during the synthesis ester/amide, ester/ester and amide/amide units can e produced in a random sequence. In order to discriminate between random and alternate ester/amide sequences MALDI-TOF/TOF-MS/MS was performed on the sodiated diamino alcohol terminated oligomers at m/z 1132 and the inspection of the data reveals the concurrent presence of some ions that allow excluding the formation of alternate ester/amide sequences. Additionally, MALDI-TOF/TOF-MS/MS was also employed to elucidate the structure of precursor ions at m/z 1115, unexpected according to the method of synthesis. Fig. 3 displays the MALDI TOF/TOF MS/MS spectrum acquired using argon as the collision gas.

Figure 3. MALDI-TOF/TOF-MS/MS spectrum and m/z values of the ions originated from the β-H transfer for the sodiated precursor ions at m/z 1114.87 Conclusions MALDI-TOF/TOF-MS/MS mass spectra showed that the main fragmentation of the PEA-Bu sample proceeds through a β-H transfer rearrangement. Some product ions proved to be diagnostic and made it possible to establish the presence of random sequences of ester and amide bonds in the PEA-Bu sample. Furthermore, MALDI-TOF/TOF-MS/MS was successfully employed to elucidate the structure of precursor ions unexpected according to the method of synthesis. References 1. P. Rizzarelli, C. Puglisi, G. Montaudo. Rapid Commun. Mass Spectrom., 19, 2407 (2005) 2. P. Rizzarelli, C. Puglisi, G. Montaudo. Rapid commun. mass spectrom., 20, 1683 (2006)

80 254 428 602 776 950m/z0

10

20

30

40

50

60

70

80

90

100

% In

tensit

y

533.39

735.52

788.57

480.32350.23

605.47

549.43278.20 392.29860.64

647.50

717.50619.45294.20 364.25 804.64122.07 747.01505.40250.17 420.32 586.51 689.56208.11

434.34

902.66178.12 331.24136.0995.05

806.63661.53551.42

462.28

HOCO(CH2)8CO-[U]2-OH


CH2 CHNHCO(CH2)8CO-[U]2-OH


CH2 CHNHCO(CH2)8COOH

CH2 CHCO-[U]2-OHCH2 CHCO-[U]1-OH

CH2 CHCO-[U]3-OH

80 254 428 602 776 95080 254 428 602 776 950m/z0

10

20

30

40

50

60

70

80

90

10

20

30

40

50

60

70

80

90

100

% In

tensit

y

533.39

735.52

788.57

480.32350.23

605.47

549.43278.20 392.29860.64

647.50

717.50619.45294.20 364.25 804.64122.07 747.01505.40250.17 420.32 586.51 689.56208.11

434.34

902.66178.12 331.24136.0995.05

806.63661.53551.42

462.28





CH2 CHNHCO(CH2)8COOH

CH2 CHCO-[U]2-OHCH2 CHCO-[U]1-OH

CH2 CHCO-[U]3-OH

CH2 CH(CH2)2NHCO(CH2)8CONH(CH2)4-OCO(CH2)8CONH(CH2)4-OCO(CH2)8CONH(CH2)4-OCO(CH2)8CONH(CH2)4OH331.2

806.6

586.4

551.4

841.6

296.2

60 262 464 666 868 1070m/z0

10

20

30

40

50

60

70

80

90

100

% In

tensit

y

586.35

788.47533.30

841.53859.56

331.19 604.35806.48

551.30

403.24

658.41349.20278.1471.65421.25 1043.66166.07 676.43463.30

445.29122.05 913.61365.21 700.47 735.37238.1589.49 190.10 620.40

256.15

303.18 505.34 990.58931.60760.56

CH2 CH(CH2)2NHCO(CH2)8COO-(CH2)4NHCO(CH2)8COO-(CH2)4NHCO(CH2)8COO-(CH2)4NHCO(CH2)8COO-(CH2)4NH2

278.2

859.6

533.4

604.4

788.5

349.2

1043.7

94.1

CH2 CH(CH2)2NHCO(CH2)8CONH(CH2)4-OCO(CH2)8CONH(CH2)4-OCO(CH2)8CONH(CH2)4-OCO(CH2)8CONH(CH2)4OH331.2

806.6

586.4

551.4

841.6

296.2

60 262 464 666 868 1070m/z0

10

20

30

40

50

60

70

80

90

100

% In

tensit

y

586.35

788.47533.30

841.53859.56

331.19 604.35806.48

551.30

403.24

658.41349.20278.1471.65421.25 1043.66166.07 676.43463.30

445.29122.05 913.61365.21 700.47 735.37238.1589.49 190.10 620.40

256.15

303.18 505.34 990.58931.60760.56

CH2 CH(CH2)2NHCO(CH2)8COO-(CH2)4NHCO(CH2)8COO-(CH2)4NHCO(CH2)8COO-(CH2)4NHCO(CH2)8COO-(CH2)4NH2

278.2

859.6

533.4

604.4

788.5

349.2

1043.7

94.1

O

OO

O

O

NH

O

O

NH

O

O

O

O

OH

O

NH

O

(1) (2)O

OO

O

O

NH

O

O

NH

O

O

O

O

OH

O

NH

O

(1) (2)

Documents

Sessione 2 Struttura, proprietà e architettura molecolare