Author`s personal copy - Laboratoire Rhéologie et Procédés

Transcription

HDR Frédéric BOSSARD
UNIVERSITE DE GRENOBLE
4
Laboratoire Rhéologie et Procédés
UMR 5520, Université Joseph Fourier, Grenoble I
Contribution à la rhéo-physique et la mise
en forme de polymères
Mémoire
présenté et soutenu publiquement le 30 novembre 2011
Pour l'obtention de l'
Habilitation à Diriger des Recherches de Grenoble Université
(Spécialité mécanique)
Par
Frédéric Bossard
Composition du jury
Rapporteurs
Examinateurs
Pr. Christian BAILLY
Pr. Christophe CHASSENIEUX
D.R. Patrick NAVARD
D.R. Nadia EL KISSI
Pr. Denis FAVIER
Pr. Patrick PIERSON
Pr. Guy SCHLATTER
IMCN – Université Catholique de Louvain
LPCI, Université du Maine
CEMEF – Mines ParisTech – Sophia Antipolis
Laboratoire Rhéologie et Procédés, UJF
Laboratoire 3SR, UJF
LTHE, UJF
LIPHT, ECPM Université de Strasbourg
A mon épouse Nathalie,
A mes enfants, Audrey, Yohann et Elouan,
A tous ceux qui m'ont permis d'apprendre et de comprendre
6
Sommaire
Chapitre 1 Curriculum Vitae détaillé
1.
Etat civil, Parcours professionnel, formation ........................................................ 4
1.1.
1.2.
1.3.
2.
Publications et communications orales ................................................................. 6
2.1.
2.2.
2.3.
2.4.
2.5.
3.
Collaborations nationales ...................................................................................................... 11
Collaborations internationales ............................................................................................... 11
Collaborations industrielles ................................................................................................... 12
Encadrement d’étudiants et jeunes chercheurs .................................................... 13
5.1.
5.2.
5.3.
5.4.
6.
Responsabilités au sein du Laboratoire ................................................................................. 10
Animation de la recherche ..................................................................................................... 10
Participation à des comités de lecture de revues internationales ........................................... 10
Collaborations et rayonnement hors laboratoire ................................................. 11
4.1.
4.2.
4.3.
5.
Revues internationales avec comité de lecture ........................................................................ 7
Chapitre d'ouvrage................................................................................................................... 8
Communications orales dans des congrès internationaux ....................................................... 8
Communications orales dans des congrès internationaux sans comité de lecture ................... 9
Communications orales dans des congrès nationaux à comité de lecture ............................. 10
Animation et management de la recherche ......................................................... 10
3.1.
3.2.
3.3.
4.
Etat civil .................................................................................................................................. 4
Parcours professionnel ............................................................................................................ 5
Formation ................................................................................................................................ 5
Masters .................................................................................................................................. 13
Projets de fin d'année............................................................................................................. 13
Thèses .................................................................................................................................... 14
Post-doctorat.......................................................................................................................... 14
Activités et responsabilités pédagogiques ............................................................ 14
6.1.
6.2.
6.3.
6.4.
En I.U.T. ................................................................................................................................ 14
En Master professionnel ........................................................................................................ 16
En formation continue ........................................................................................................... 16
A l'international ..................................................................................................................... 16
Chapitre 2 Activités de recherche
1.
Résumé des activités de recherche ....................................................................... 18
2.
Synthèse des principaux résultats ........................................................................ 23
2.1. Rhéo-physique de polymères associatifs ...................................................................... 23
2.2. Mise en forme et caractérisation de composites .......................................................... 40
Chapitre 3 Perspectives
Annexes ..................................................................................................................... 57
Remerciements
Ce manuscrit retrace mon activité de recherche réalisée depuis ma thèse, soutenue en
2001. Ce travail n'aurait pu être mené sans une collaboration étroite avec plusieurs personnes que
je souhaite remercier.
Je souhaite tout d'abord exprimer ma profonde gratitude aux professeurs Thierry AUBRY
et Michel MOAN de l'équipe rhéologie du laboratoire d'ingénierie des matériaux de Bretagne
(LIMATB), de Brest. Ils ont su me transmettre leur passion pour la recherche, ils m'ont accordé
leur confiance et m'ont donné de précieux conseils. Je garde en mémoire leurs soucis constants
de rigueur et d’exactitude et je leur dois ce que je suis professionnellement.
Je remercie chaleureusement Dr. Spyros YANNOPOULOS, Pr. Georgios STAIKOS et
Pr. Constantinos TSITSILIANIS de l'Institute of Chemical Engineering (ICE-HT/FORTH) de
Patras, Grèce, avec qui j'ai collaboré pendant mon post-doctorat. Ils ont fait preuve d’une
hospitalité que je n’oublierai pas et m’ont permis de travailler dans d'excellentes conditions.
A Albert MAGNIN, directeur du Laboratoire Rhéologie et Procédés de l’Université
Joseph Fourier qui m’a accueilli comme jeune maître de conférences. Je tiens à le remercier pour
la politique de recherche qu'il a mené au laboratoire, visant à faciliter l'intégration et le travail des
jeunes chercheurs du laboratoire.
A mon arrivée au Laboratoire de Rhéologie, j’ai intégré l’équipe Polymères et Mise en
Œuvre, dirigée par Nadia EL KISSI. Je tiens à lui exprimer toute ma reconnaissance pour la
confiance qu'elle m'a accordée, son écoute et sa disponibilité.
Mon activité de recherche expérimentale a été rendu possible par l'appui technique du
Laboratoire Rhéologie et Procédés par l'intermédiaire d'Hélène GAILLARD, Didier BLESES,
Frédéric HUGENELL, Mohammed KARROUCH et Eric FAIVRE sans oublier le travail du
secrétariat mené par François BERGEROT, Claudine LY-LAP et Sylvie GAROFALO.
Je remercie chaleureusement mes collègues et amis du Laboratoire Rhéologie et Procédés
pour la qualité de l'ambiance de travail qu'ils ont su créer. Je remercie particulièrement JeanRobert CLERMONT pour les remarques et conseils qu’il m’a donnés dans le cadre de l’examen
de ce mémoire.
La plupart des travaux présentés ici n’auraient jamais été développés sans les doctorants
que j’ai eu la chance de co-encadrer. Je remercie ainsi Alessendra D'APREA et Anica
LANCUSKI ainsi que Bibekananda SUNDARY en post-doctorat au laboratoire.
Je remercie enfin Messieurs Christian BAILLY, professeur à Institut de la Matière
Condensée et des Nanosciences, Université Catholique de Louvain, Christophe
CHASSENIEUX, professeur au Laboratoire Polymères, Colloïdes, Interfaces, Université du
Maine, et Patrick NAVARD, Directeur de recherche CEMEF – MINES ParisTech, Sophia
Antipolis, pour avoir accepté d’être rapporteurs de ce mémoire et membres du jury.
Je remercie Nadia EL KISSI, directrice de recherche au Laboratoire Rhéologie et
procédés, Denis FAVIER, Professeur au Laboratoire Sols, Solides, Structures et Risques,
Université Joseph Fourier, Patrick PERSSON, Professeur au Laboratoire d'Etude des Transferts
en Hydrologie et Environnement et Guy SCHLATTER, professeur au Laboratoire d’Ingénierie
des Polymères pour les Hautes Technologies, ECPM Université de Strasbourg, qui ont accepté
de participer à ce jury.
8
Chapitre 1
4
Curriculum Vitae détaillé
1. Etat civil, Parcours professionnel, formation
1.1. Etat civil
Prénom Nom :
Frédéric BOSSARD
Date de naissance :
Lieu de naissance :
Situation familiale :
19 décembre 1970
Landivisiau (29), Finistère
Marié, 3 enfants
Adresse professionnelle :
Laboratoire de Rhéologie
1301 rue de la Piscine
Domaine Universitaire
BP 53 - 38041 Grenoble cedex 9
Adresse personnelle :
748 Grande Rue
38660 Le Touvet
[email protected]
http://rheologie.ujf-grenoble.fr/
(33) 4 76 82 51 79
(33) 4 76 82 51 64
Adresse e-mail :
Site internet:
Téléphone professionnel :
Fax :
IUT 1 Grenoble, Département GMP
151, rue de la Papeterie
BP 67
38402 Saint Martin D'Hères
5
Etat civil, Parcours professionnel, formation
1.2. Parcours professionnel
Sept. 2006
Maître de conférences à l'I.U.T 1 de Grenoble, Département
GMP; Laboratoire de rhéologie, UJF, UMR5520, Grenoble INP
Janv. 2004 - Déc. 2005
Post-doctorat - Laboratoire Polymères, Propriétés aux Interfaces
et Composites, (L2PIC) - Lorient
Oct. 2002 - Oct. 2003
Post-doctorat - Institute of Chemical Engineering and High
Temperature Chemical Processes, (ICE-HT/FORTH) - Patras,
Grèce
Sept. 2000 - Sept. 2002
Attaché temporaire d'enseignement et de recherche Université de Bretagne Occidentale (U.B.O.).
1.3. Formation
1998 – 2001
Doctorat en Physique – Université de Bretagne Occidentale (U.B.O.),
Brest.
Sujet de Thèse : "Etude rhéologique de suspensions aqueuses diluées et
concentrées de plaquettes d'argile colloïdales. Effets de l'adsorption de
polymères associatifs." Mention très honorable
Directeur de Thèse : Michel Moan.
Jury : M. Tassin (Université du Maine)
M. Van Damme (ESPCI)
M. Aubry, (U.B.O.)
Mme Audibert-Hayet (Institut Français du Pétrole)
M. Moan (U.B.O.)
1994 – 1997
Elève Officier Pilote de l'Aéronautique Navale.
1993 – 1994
D.E.A. de Physique, option matière condensée et matériaux Université de Rennes I. Mémoire : "Propriétés rhéologiques de suspensions
d'argile synthétique"
1992 – 1993
Maîtrise de Physique - Université de Rennes I.
1991 – 1992
Licence de Physique - Université de Rennes I
1989 – 1991
D.E.U.G. A, mathématiques, physique et informatique - U.B.O.
6
2. Publications et communications orales
La diffusion de mon travail de recherche s'est traduite par la rédaction de 17 articles dans des
revues à comité de lecture, 1 chapitre d'ouvrage, 17 communications orales dans des congrès
internationaux, 3 communications orales dans des congrès nationaux à comité de lecture.
Type de publications
Articles dans des revues
internationales
Chapitre d'ouvrage
Titre de la revue
Journal of Rheology
4
Langmuir
2
Macromolecules
2
Rheologica Acta
Soft Matter
Polymer
Polymer Engineering & Science
Polymeric Materials: Science and Engineering
Physical Review Letters
SPE journal
Electrochimica Acta
Cellulose
Hydrogen-Bonded Interpolymer Complexes: Formation,
Structure and Applications
1
1
1
1
1
1
1
1
1
Année
2003
2003
2004
2007
2002
2007
2004
2005
2010
2006
2002
2008
2004
2004
2003
2010
2011
1
2009
Annual European Rheology Congress
4
2011
2010
2006
2002
International Congress on Rheology
2
2008
1
1
1
1
2009
2001
2004
2001
1
2005
1
1
1
2005
2005
2000
1
2002
1
1
2004
2003
2003
2010
2007
de Gennes discussion Conference
International Meeting of the Hellenic Society of Rheology
Annual Congress of the Hellenic Society of Rheology
Pacific Rim Conference on Rheology
Congrès internationaux
International Symposium on Nanostructured and Functional
Polymer-based Materials and Composites
European Polymer Congress
ACS Colloid and Surface Science Symposium
Canadian Chemical Engineering Conference
Society of Petroleum Engineers Annual Technical Conference
and Exhibition
American Chemical Society National Meeting
Workshop
Congrès nationaux
Nbre
Congrès annuel du Groupe Français de Rhéologie
2
Congrès Français de Mécanique
1
7
Publications et communications orales
Une présentation chronologique et détaillée des publications est proposée ci-dessous:
2.1. Revues internationales avec comité de lecture
[1] Alloin F. , A. D’Aprea, A. Dufresne, N. El Kissi, F. Bossard, "Poly(oxyethylene) and
ramie whiskers based nanocomposites. Influence of processing: extrusion and
casting/evaporation", Cellulose, 18, 957 – 973, 2011.
Taux de citation: 0
[2] Alloin F. , A. D’Aprea, A. Dufresne, N. El Kissi, F. Bossard, " Nanocomposite polymer
electrolyte based on whisker or microfibrils polyoxyethylene nanocomposites",
Electrochimica Acta, 55, 5186–5194, 2010.
Taux de citation: 0
[3] Bossard, F., N. El Kissi, A. D'Aprea, F. Alloin, J-Y Sanchez and A. Dufresne, " Influence
of dispersion procedure on rheological properties of aqueous solutions of high molecular
weight PEO", Rheologica Acta. 49, 529–540, 2010.
Taux de citation: 4
[4] Bossard, F., I. Pillin, T. Aubry and Y. Grohens, "Rheological Characterization of Starch
Derivatives/Polycaprolactone Blends Processed by Reactive Extrusion", Polym. Eng. Sci,
48, 1862-1870, 2008.
Taux de citation: 1
[5] Sotiropoulou, M, F. Bossard, E. Balnois, J. Oberdisse and G. Staikos, "Characterization of
the Core-Shell Nanoparticles Formed as Soluble at low pH Hydrogen bonding
Interpolymer Complexes", Langmuir, 23, 11252 –11258, 2007.
Taux de citation: 4
[6] Bossard, F., M. Moan, T. Aubry, "Linear and non-linear Viscoelastic Behavior of Very
Concentrated Kaolinite Suspensions", J. Rheol. 51, 1253-1270, 2007.
Taux de citation: 7
[7] Bossard, F., T. Aubry, G. Gotzamanis, C. Tsitsilianis, "pH-tunable Rheological Properties
of a Telechelic Cationic Polyelectrolyte Reversible Hydrogel" Soft Matter, 2, 510-516, 2006.
Taux de citation: 32
[8] Bossard, F., V. Sfika, C. Tsitsilianis and S. Yannopoulos, "A Novel Thermothickening
Phenomenon Exhibited by a Triblock Polyampholyte in Aqueous Salt-Free Solutions",
Macromolecules, 38, 2883-2888, 2005.
[9] Bossard, F., M. Sotiropoulou and G. Staikos, "Thickening effect in Soluble Hydrogenbonding Interpolymer complexes. Influence of molecular composition and pH" J. Rheol.,
48(4), 927-926, 2004.
[10] Tsitsilianis, C., F. Bossard , V. Sfika, N. Stavrouli, A. Kiriy, G. Gorodyska, M. Stamm, and
S. Minko, "Multifunctional Double Hydrophilic Triblock Copolymer in Solution and on
Surface." Polym. Mater. Sci .& Engineering, 90, 368-369, 2004.
Taux de citation: 0
[11] Scopigno T., R. DiLeonardo, G. Ruocco, A.Q.R. Baron, S. Tsutsui, F. Bossard, S.N.
Yannopoulos, "High frequency dynamics in a monatomic glass", Phys. Rev. Lett, 92(2),
025503, 2004.
[12] Bossard, F., V. Sfika, C. Tsitsilianis "Rheological Properties of Physical Gel formed by
Triblock Polyampholyte in Salt-Free Aqueous Solutions", Macromolecules, 37, 3899-3904,
2004.
[13] Herzhaft, B., L. Rousseau, L. Neau, M. Moan and F. Bossard, "Influence of temperature
and clays/emulsion microstructure on oil-based mud low shear rate rheology", SPE
Journal, 8 (3): 211-217, 2003.
Taux de citation: 5
[14] Moan, M., T. Aubry, F. Bossard, "Nonlinear Behavior of Very Concentrated Suspensions
of Plat-like Kaolin Particles in Shear Flow" J. Rheol., 47(6); 1493-1504, 2003.
[15] Aubry, T., F. Bossard, G. Staikos, G. Bokias, "Rheological study of semi-dilute aqueous
solutions of a thermoassociative copolymer", J. Rheol., 47(2); 577-587, 2003.
[16] Aubry, T., F. Bossard and M. Moan, "Rheological Study of Compositional Heterogeneity in
an Associative Commercial Polymer Solution", Polymer, 43; 3375-3380, 2002.
Taux de citation: 5
[17] Aubry, T., F. Bossard and M. Moan, "Laponite Dispersions in the Presence of an
Associative Polymer.", Langmuir, 18(1); 155-159, 2002.
2.2. Chapitre d'ouvrage
[1] Staikos G., M. Sotiropoulou, G. Bokias, F. Bossard, J. Oberdisse, E. Balnois, Chapitre 2,
"Hydrogen-Bonded Interpolymer Complexes Soluble at Low pH" publié dans " HydrogenBonded Interpolymer Complexes: Formation, Structure and Applications", editeur: World
Scientific Publishing Co., mars 2009., ISBN 978-981-270-785-7
2.3. Communications orales dans des congrès internationaux
[1]
[2]
[3]
[4]
[5]
[6]
Bossard F., Sundaray, B., Lancuski, A. and Pétrier, C. "Influence of elongational properties
of polymer solutions on nanofibre properties processed by electrospinning", 6th Annual
European Rheology Conference, Suzdal, Russie, 2011.
Bossard F., El Kissi N, D'Aprea A, Alloin F, Sanchez J-Y and Dufresne A "Rheological
investigation of polymer scission and aggregation induced by the dispersion in high
molecular weight PEO solutions", 6th Annual European Rheology Conference, Göteborg,
Suède, 2010.
Bossard F., El Kissi N, D'Aprea A, Alloin F, Sanchez J-Y and Dufresne A, "Influence of
turbulent flow on rheological properties of aqueous solutions of high molecular weight
PEO", de Gennes Discussion Conference, Chamonix, France, 2009.
El Kissi N, F. Alloin, A. Dufresne, F. Bossard and A. D'Aprea, "Influence of cellulose
nanofillers on the rheological properties of polymer electrolytes", 15th International
Congress on Rheology/80th Annual Meeting of the Society-of-Rheology, Monterey, CA.,
American Institute of Physics Conference Proceedings, 1027, 87-89, 2008.
Bossard F., M. Moan and T. Aubry, "Very concentrated plate-like kaolin suspensions
under large amplitude oscillatory shear: A microstructural approach", 15th International
Congress on Rheology/80th Annual Meeting of the Society-of-Rheology, Monterey, CA.,
American Institute of Physics Conference Proceedings, 1027, 695-697, 2008.
Aubry T., F. Bossard, G. Gotzamanis and C. Tsitsilianis, "Rheological Properties of
Cationic Telechelic Polyelectrolyte", 3rd Annual European Rheology Conference,
Hersonisos, Grèce, 2006.
8
9
Publications et communications orales
[7]
[8]
[9]
[10]
[11]
[12]
[13]
[14]
[15]
[16]
[17]
Tsitsilianis, C., I. Katsampas, F. Bossard., V. Sfika, A. Kiriy, G. Gorodyska, M. Stamm, S.
Minko "Responsive triblock copolymers with amphoteric blocks: from ordered to aperiodic
nanostructures", 1st International Symposium on Nanostructured and Functional Polymerbased Materials and Composites, Dresden, Allemagne, 2005.
Gotzamanis, G, F. Bossard, R. Lupytsky, T. Aubry, S. Minko, C. Tsitsilianis, "Physical
gelation and rheological properties of cationic telechelic polyelectrolytes", 79th ACS Colloid
and surface Science symposium, Potsdam, NY, 2005.
Sotiropoulou M., F. Bossard, Bokias G., Oberdisse J., Staikos G., "Soluble Hydrogenbonding Interpolymer Complexes and their pH controlled Gel-like behaviour in Water",
European Polymer Congress, Moscou, Russie, 2005.
Tsitsilianis C., F. Bossard, Stavrouli N. and Sfika V.,"The Rheology of a Physical Gel
Formed by Double Hydrophilic Triblock Copolymers", Annual Congress of the Hellenic
Society of Rheology, Athens, Greece, 2004.
Tsitsilianis, C., F. Bossard, V. Sfika, N. Stavrouli, A. Kiriy, G. Gorodyska, M. Stamm and
S. Minko, “Multifunctional Double Hydrophilic Triblock Copolymer in Solution and on
Surface”, 227th American Chemical Society National Meeting, Anaheim, Californie, 2004.
Sotiropoulou, M., F. Bossard, C. Cincu, G. Bokias and G. Staikos, "Water-soluble
polyelectrolyte and hydrogen-bonding interpolymer complexes for nanoparticles
formation", EPF, Nanostructured Polymer Materials, Gargnano, Italy, 2003.
Aubry, T., F. Bossard, M. Moan, "Rheological Properties of a Thermoassociative polymer
in Aqueous Solutions.”, 6th European Congress on Rheology, Erlangen, Allemagne, 2002.
Herzhaft, B., L. Rousseau, L Neau, M. Moan, F. Bossard, "Influence of temperature and
clays/emulsion microstructure on oil based muds low shear rate rheology." Society of
Petroleum Engineers Annual Technical Conference and Exhibition, San Antonio, USA,
2002.
Aubry, T., F. Bossard, M. Moan, "Influence of macromolecular associations on the dilute
properties of a clay suspension.", 3rd International Meeting of the Hellenic Society of
Rheology, Patras, Grèce, 2001.
Moan, M., T. Aubry, F. Bossard, "Nonlinear behavior of concentrated suspensions of
platelike kaolinite particles in shear flow", 3rd Pacific Rim Conference on Rheology,
Vancouver, B.C., Canada, 2001.
Moan, M., T. Aubry, F. Bossard, "Etude de quelques propriétés non-linéaires de
suspensions concentrées de kaolin.", 50th Canadian Chemical Engineering Conference,
Montreal, Canada, 2000.
2.4. Communications orales dans des congrès internationaux sans comité de lecture
[1]
[2]
[3]
[4]
Bossard F., El Kissi N, D'Aprea A, Alloin F, Sanchez J-Y and Dufresne A, "The
unsuspected brittleness of high molecular weight PEO in solution: influence of the
dispersion procedure", Rheology Alpine Meeting, Tignes, 2011.
Bossard F, "A review on polymer nanofibers processed by electrospinning", Rheology
Alpine Meeting, Les Saisies, 2010.
Bossard F, Aubry T., M. Moan "Rheological behaviour of very concentrated suspensions of
plate-like clay particles", Rheology Alpine Meeting, Chatel, 2008.
Bossard F. Tsitsilianis, C. "Influence of concentration and temperature on rheological
properties of a novel water-soluble polyampholyte", Rheology Alpine Meeting, Samoëns,
2007.
2.5. Communications orales dans des congrès nationaux à comité de lecture
Bossard F., El Kissi N, D'Aprea A, Alloin F, Sanchez J-Y and Dufresne A, "Elaboration
et mise en forme de polymères nano-composites pour batteries au lithium", 45ème congrès
annuel du Groupe Français de Rhéologie, Lyon, France, 2010.
Bossard F., M. Moan et T. Aubry, "Comportement viscoélastique linéaire et faiblement
non-linéaire de suspensions concentrées de plaquettes d'argile colloïdales" 18ème congrès
Français de Mécanique, Grenoble, France, 2007.
Bossard, F., C. Tsitsilianis et S. Yannopoulos, "Etude rhéologique de solutions d'un
polyampholyte tribloc: effet du précisaillement et de la température", 38ème congrès annuel
du Groupe Français de Rhéologie, Brest, France, 2003.
[1]
[2]
[3]
3. Animation et management de la recherche
3.1. Responsabilités au sein du Laboratoire
-
Représentant élu du personnel de catégorie B (chargé de recherche et maître de
conférences)
Responsable de l'organisation du congrès international "Annual Alpine Rheology Meeting"
depuis 2011.
Responsable de la Formation Continue
3.2. Animation de la recherche
-
-
-
-
Responsable du projet Nano4neuro du pôle SMingue de l'UJF (2010 – 2011)
"Mise en forme d'une structure de nanofibres biopolymère par electrospinning
dédiée aux neurosciences"
Partenaire : F. Mc Cluskey, LEGI, C. Petrier, LEPMI
Financement: 30 000 Euros
Partenaire dans le Bonus Qualité Recherche de Grenoble INP
"Nouvelles batteries lithium mettant en jeu des matériaux d’électrode et séparateur
organiques innovants."
Responsable: Jean-Claude Leprêtre, LEPMI
Financement: 1 an de post-doctorat et 10 000 Euros
Membre élu du conseil du Groupe Français de Rhéologie (GFR) depuis novembre 2009.
Responsable du site du GFR
Chairman du 6ème congrès annuel européen de rhéologie, cession "melt polymers and
solutions", 8 avril 2010, Gôteborg, Suède.
Rapporteur de la thèse de Mr. Shri. K P Rajesh; "Studies of Ion Conducting Properties of
Electrospun Polymer Fibers", Department of Physics, Indian Institute of Technology
Madras, Inde.
Membre de 2 comités de sélection (postes 62/60McF1131 en 2010 et 60McF0419 en
2011)
3.3. Participation à des comités de lecture de revues internationales
Octobre
Février
Juillet
2006
2007
2007
Soft Matter
Physical Chemistry Chemical Physics
Mechanics of Materials,
10
11
Collaborations et rayonnement hors laboratoire
Septembre
Janvier
Decembre
Juin
Novembre
Mars
Mars
2007
2008
2008
2009
2009
2010
2011
Journal of Chemical Physics,
Polymer Engineering and Science,
Polymer Engineering and Science,
Colloids and Surfaces A: Physicochemical and Engineering Aspects,
Colloids and Surfaces A: Physicochemical and Engineering Aspects
Polymer Engineering and Science
Rheologica Acta
4. Collaborations et rayonnement hors laboratoire
4.1. Collaborations nationales
Mon parcours professionnel m'a permis de nouer des collaborations à l'échelle nationale avec
différents laboratoires et organismes de recherche.
- Entre le Laboratoire des Polymères, Propriétés aux Interfaces et Composites, L2PIC,
Lorient et le laboratoire de Rhéologie de Brest, dans le cadre de mon post-doctorat
Thématique: Mise en forme et caractérisation de bio-composites de type PCL chargé en
amidon.
- Entre le Laboratoire de Rhéologie de Grenoble et:
o Le Laboratoire d'Electrochimie et de Physicochimie des Matériaux et des
Interfaces, LEPMI et le Laboratoire Génie des Procédés Papetiers, LGP2 dans le
cadre de la thèse d'Alessandra D'Apréa et le post-doctorat de Bibekananda
Sundaray.
Thématique: Mise en forme et caractérisation de membranes polymère pour piles
Lithium
o Le Centre de Recherche sur les Macromolécules Végétales, CERMAV et l'équipe
Régénération et Croissance de l'Axone du Laboratoire Physiopathologie des
Maladies du Système nerveux Central, INSERM UMRS-952, CNRS UMR 7224,
Université Pierre et Marie Curie, Paris VI, dans le cadre de la Thèse d'Anica
Lancuski
Thématique : Elaboration et caractérisation de nanofibres de bio-polymères
fonctionnalisés pour des applications en neurosciences
o L'équipe Mécanique et Couplages Multiphysiques des Milieux Hétérogènes
(CoMHet) du Laboratoire Sols Solides Structures Risques, 3S-R pour la
caractérisation mécanique et tomographie X des structures de fibres de polymères
obtenues par electrospinning
4.2. Collaborations internationales
J'ai eu la volonté, dès la fin de ma thèse, de développer des collaborations internationales.
- Entre l'Institute of Chemical Engineering de Patras, ICE-HT/FORTH et le Department of
Chemical Engineering de l'université de Patras en Grèce pendant mon poste doctorat en
2002 – 2003. Ce poste m'a permis de collaborer avec des spécialistes de la synthèse des
polymères en la personne de Georgios Staikos et Constantinos Tsitsilianis sur des systèmes
de type polyampholyte, téléchélique et des mélanges de polymères complexant.
- Entre la fin juin et la fin aout 2009, j'ai été invité par le professeur Pierre Carreau, directeur
du Centre de Recherche en Plasturgie et Composite, (CREPEC) et professeur à l'Ecole
Polytechnique de Montréal. A cette occasion, j'ai débuté des travaux portant sur le
développement de polymères biodégradables à base de nanocristaux de cellulose. J'ai, par
ailleurs suivi les travaux de thèses de trois étudiants du laboratoire, l'un portant sur des
résines chargées en nanotubes de carbone, le second sur des mousses de thermoplastiques et
le troisième sur la mise en forme de nanofibres par electrospinning.
- Je participe depuis 2009 à la construction d'un réseau d'expertise et de formation en
rhéologie élongationnelle. Ce réseau regroupe les membres du conseil de l'European Society
of Rheology, ESR, avec Mats Stading (The Swedish Institute for Food and Biotechnology,
Suède et porteur du projet), Mike Webster (The Institute of Non-Newtonian Fluid
Mechanics, Swansea University, Grande-Bretagne), Igor Emri (université de Ljubliana,
Slovénie), Críspulo Gallegos (université d’Huelva, Espagne), Joao Maia (Case Western
Reserve University, USA). Cette collaboration s’est traduite par la rédaction de projets de
recherche européens FP7-PEOPLE-ITN en 2008 intitulé "Rheoextens", 2010
"Rheocompetence" et 2011 "Rheotraining" dont j'étais le coordinateur local pour le
laboratoire de Rhéologie. Ces projets de recherche ont pour but d'établir un réseau
d'excellence entre ces différents instituts et des industriels autour de la rhéologie
élongationnelle, proposant aux jeunes chercheurs européens des cours de rhéologie, des
séminaires, des conférences et d'effectuer une partie de leurs travaux dans un laboratoire
membre du réseau.
4.3. Collaborations industrielles
2007-2008 : Contrat avec la société Maurel & Prom (contrat n° 522070026 INPG entreprise
SA, montant 76000 Euro). Cette société française spécialisée dans la récupération assistée du
pétrole nous a contactés, par l'intermédiaire de Mr. Philippe Royer, ingénieur réservoir, pour
répondre aux problèmes suivants: caractériser les propriétés rhéologiques de leurs bruts
paraffiniques additivés en fonction de la température et concevoir et réaliser un rhéomètre
destiné aux lieus de production. L'enjeu industriel pour la société Maurel & Prom était :
- de sélectionner un additif permettant d'abaisser la température d'écoulement du
pétrole
- de dimensionner leurs installations de pompage (puissance des pompes et nombre de
réchauffeurs) en déterminant les contraintes de redémarrage des installations et en
étudiant l'influence des cycles thermiques sur les propriétés rhéologiques des bruts.
- de pouvoir caractériser et ajuster rapidement la viscosité de leurs bruts.
J'ai participé à l'élaboration du devis lors du montage du projet. En collaboration avec Laurent
Jossic du Laboratoire de Rhéologie de Grenoble, nous étions en charge de la caractérisation
rhéologique des fluides. Ce projet a fait l'objet de trois présentations scientifiques auprès
d'industriels (Total, Scomi Anticor, Arkema, Roemex).
2009-2010 : Collaboration avec la société CHIMEC S.p.A. Cette multinationale italienne est
spécialisée dans la formulation d'additifs dédiés à l'exploitation pétrolière et l'industrie de
raffinage pour résoudre les problèmes liés à la corrosion, aux dépôts et aux phénomènes
d'émulsion inverse. J'ai été invité au siège de la société CHIMEC à Santa Paloma en octobre 2009
par Mr. Marco Romagnoli et Mr. Marcello Della Corte, responsable production Afrique pour
définir des possibilités de collaborations scientifiques. Dans un premier temps, nous avons
convenu de former aux techniques rhéométriques pendant 15 jours des ingénieurs de recherche
de la société. Cette formation est programmée en fin d'année 2011. Par ailleurs CHIMEC
souhaite étendre les collaborations dans le cadre de stages d'étudiants grenoblois au Laboratoire
de Rhéologie et dans leur centre de R&D.
12
13
Encadrement d’étudiants et jeunes chercheurs
5. Encadrement d’étudiants et jeunes chercheurs
5.1. Masters
Mai - Juillet 2011
Mai - Juillet 2011
Février - Juin 2011
Avril - Juillet 2009
Mars - juillet 2009
Avril - Juillet 2008
Mars - Juin 2007
Michaël Chiarappa, 4ème année Polytech'Annecy, spécialité Matériaux
composites.
Sujet: Etude du phénomène de contraction de nanofibres de PVdF
mises en forme par electrospinning
Taux d'encadrement : 100%
Sébastien Dagaz, 4ème année Polytech'Annecy, spécialité Matériaux
composites.
Sujet: Elaboration et caractérisation de nanofibres de polystyrènes:
condition d'obtention de fibres non poreuses.
Jeferson de Olivera, Master en Mechanical Engineering, Federal
University of Minas Gerais, Bresil
Sujet: Design and development of a rotative collector for electrospinning
Meriem Abdelkhalek, Master en Gestion des Systèmes Industriels,
spécialité Formulation, Analyse et Contrôle.
Sujet: Caractérisation rhéologique de solutions de polymères
associatifs hydrophobes: Influence du pH et de la concentration
Taghrid Mhalla, Master en Gestion des Systèmes Industriels, spécialité
Formulation, Analyse et Contrôle, étudiante Syrienne.
Sujet: Dynamique moléculaire des systèmes auto associatifs
Taux d'encadrement : 50% (co-encadrement avec Yahia Rharbi, Chargé
de Recherche CNRS, Laboratoire de Rhéologie)
Nicolas Manu, Master en Gestion des Systèmes Industriels, spécialité
Formulation, Analyse et Contrôle
Sujet: Dynamique moléculaire de l'échelle micrométrique à
l'échelle macroscopique de solutions de polymères associatifs: Les
HASE
Taux d'encadrement : 50% (co-encadrement avec Yahia Rharbi, Chargé
Moulay Abdelkarim Elmoussaoui, Master à UFR de Mécanique de
Grenoble
Sujet: Elaboration et mise en forme de polymères nanocomposites
Taux d'encadrement : 50% (co-encadrement avec Nadia El Kissi, Chargé
5.2. Projets de fin d'année
Février-Juillet 2007
Olivier Pras, Ecole Nationale Supérieure d'Hydraulique et de Mécanique
de Grenoble
Sujet: Caractérisation rhéologique d'un pétrole brut paraffinique
Contrat industriel avec la société Maurel & Prom (montant de 76 000 €)
Taux d'encadrement : 50% (co-encadrement avec Laurent Jossic, Maître
de conférences, Laboratoire de Rhéologie)
5.3. Thèses
Sept 2010 -
Anica Lancuski (nationalité serbe)
Sujet: Mise en forme de nanofibres de bio-polymères fonctionnels
par electrospinning pour des applications en neurosciences
Taux d'encadrement : 50% (co-encadrement avec Sébastien Fort, Chargé
de Recherche CNRS, CERMAV)
1 présentation orale au 7th Annual European Rheology Conference,
Suzdal, Russie 2011.
Sept 2005- Juin 2009 Alessandra D’Aprea, (nationalité italienne)
Sujet: Etude rhéologique et physico-chimique de membranes
nanocomposites PEO/cellulose pour batterie au Lithium.
Influence du procédé d'élaboration et de la nature des
nanoparticules de cellulose
Taux d'encadrement : 33% (co-encadrement avec Nadia El Kissi,
Directeur de recherche CNRS, Laboratoire de Rhéologie; Fannie Alloin,
Directeur de recherche CNRS, Laboratoire d'Electrochimie et de
Physicochimie des Matériaux et des Interfaces, LEPMI)
3 articles publiés
Devenir du docteur : Post-doctorat au CEA de Grenoble
5.4. Post-doctorat
Octobre 2010
Bibekananda Sundaray (nationalité indienne)
Sujet: Nouvelles batteries lithium mettant en jeu des matériaux
d’électrode et séparateur organiques innovants
6. Activités et responsabilités pédagogiques
L'enseignement et les responsabilités administratives nécessaires au bon fonctionnement
des composantes pédagogiques représentent une part importante de l'activité d'un enseignant
chercheur. Je retrace dans ce chapitre mes principales implications dans ces domaines.
6.1. En I.U.T.
J'ai été nommé maître de conférences en 2006 au département Génie Mécanique et
Productique, (GMP) de l'IUT 1 de Grenoble. Les besoins en termes d'enseignement et
d'encadrement sont forts dans les IUT. Afin de faciliter le bon fonctionnement du département
GMP et mon intégration au sein de l'équipe pédagogique, j'ai décliné la possibilité de décharge
d'enseignement accordée par l'Université Joseph Fourier à tout maître de conférences
nouvellement nommé.
Je suis responsable de l'enseignement de la mécanique en 1ère année: je coordonne les
enseignants intervenant dans les différents modules de la mécanique, le contenu pédagogique et
d'évaluation des connaissances. Je suis également responsable du laboratoire COMETHE
(COnception, MEcanique et THErmique) depuis ma nomination: ce laboratoire est destiné
principalement aux travaux pratiques des étudiants du département GMP. J'ai renouvelé et
augmenté l'éventail des TP de mécanique et de dimensionnement des structures (DDS) pour un
budget total de 23 500 euros.
14
15
Activités et responsabilités pédagogiques
Mes enseignements dans cette discipline se déclinent comme suit:
Modules
Statique – F113
Cinématique – F 214
Cinétique – F215
Dynamique – F313
Energétique – F 314
DDS – F 312
Total
Heures de
cours
-10
8
Heures de TD Heures de TP
18
40
32
----
--
--
18
18
90
6
24
Tableau 1 : Répartition des enseignements en mécanique
J'interviens également en mathématiques en 1ère et 2ème année comme suit:
Modules
Dérivée, différentielles – F111
Calcul intégral – F 115
Fonction à plusieurs variables
– F212
Calcul matriciel – F311
Courbes – F 411
Total
Heures de cours
--9
(Depuis 2009)
--9
Heures de TD
18
10
32
18
(sauf en 2009)
10
88
Tableau 2 : Répartition des enseignements en mathématiques
Entre 2008 et 2010, j'ai accepté la responsabilité des projets de 1ère année au département
GMP de l’IUT1. Par binômes, les étudiants choisissent un mécanisme simple qu'ils étudient tout
au long de l'année. Au premier semestre, une étude cinématique et une description des solutions
technologiques retenues pour la réalisation des mécanismes leur sont demandées. Au second
semestre, une étude plus approfondie est réalisée, comprenant le cahier des charges fonctionnel,
les modes d'obtention des pièces, une analyse expérimentale des matériaux utilisés. Ces projets
font l'objet de présentations orales en français et en anglais et de rédaction d'un rapport
technique. Outre leur importance pédagogique, ces projets représentent pour chaque semestre
une unité d'enseignement dont l'évaluation s'avère décisif pour le passage en 2ème année. En tant
que responsable de ces projets, j'avais un travail d'information, de communication et d'écoute
auprès des étudiants, d'organisation des suivis de projets avec les collègues enseignants en
construction, fabrication, sciences des matériaux, anglais et de l'évaluation.
Depuis 2010, je suis responsable de la poursuite d'études des étudiants de 2 ème année du
département GMP. Bien que la vocation première des IUT est de former en deux ans des
techniciens supérieurs, près de 90% de nos étudiants poursuivent leur formation en France ou à
l'étranger. J'organise un jury de poursuite d'étude en fin de semestre 3 au cours duquel un avis de
poursuite d'étude est notifié à chaque étudiant. Je reçois en entretien individuel les étudiants de
2ème année pour définir leur projet de poursuite d'étude et les conseiller. Je veille à la préparation
de chaque dossier (près de 350 dossiers par an) en collaboration avec Mme Ardid de
l'administration centrale de l'IUT 1 de Grenoble.
J'ai été élu au conseil de département pour représenter les enseignants du supérieur au
sein du département GMP.
La formation en IUT accorde une place importante aux stages en entreprise. J'encadre
chaque année entre 4 à 6 étudiants de 2ème année de GMP et un étudiant de licence
professionnelle en alternance.
6.2. En Master professionnel
Depuis septembre 2007, j'assure la formation en rhéologie des fluides complexes en Master 1,
Chimie et Procédés Majeurs, Génie des Systèmes Industriels, option Formulation, Analyse,
Contrôle à l'UFR de chimie de l'université Joseph Fourier. Cette formation comprend 33 heures
de cours/TD.
Enfin, depuis la rentrée 2011, je suis responsable de l’Unité d'Enseignement "Matière
divisée et suspension" en Master 1, Génie des systèmes Industriels. Cette UE regroupe, outre la
rhéologie, un module de physicochimie et interfaces et un module d'agitation et mélange.
6.3. En formation continue
Dans le cadre de la formation continue de Grenoble Institut National Polytechnique
(Grenoble INP), je participe à la formation d'industriels en rhéologie.
Avril 2008: Société ABB spécialisée dans la conception et la réalisation d'installation de
formulation de peinture
Publique: 12 ingénieurs
Durée: 14 heures soit 2 jours sur les 3 journées de formation
Juin 2008:
Formation intra-Grenoble INPG
Publique: 4 ingénieurs (1 de Nestlé Bauvais, 2 de MicroChemical en Suisse, et 1 de
AXENS usine de Salindres)
Durée: 10 heures soit 1 jour et demi sur les 3 journées de formation
Juin 2009:
Société Bayer CropScience leader dans le marché de la protection des cultures
Publique: 5 ingénieurs
Durée: 14 heures sur 2 jours
Décembre 2010:
Société Théramex SAM, filiale du groupe pharmaceutique Merck spécialisée
dans la pharmacologie gynécologique
Publique : 15 ingénieurs (galénistes, pharmaciens, formulateurs)
Durée: 14 heures sur 2 jours
6.4. A l'international
Dans le cadre du programme européen ERASMUS et sur le principe d’un accord bilatéral
entre l'IUT 1 de Grenoble et le Cork Institute of Technology (CIT), j'ai participé à une mission
"Teaching Staff" au cours de l'année universitaire 2008-2009. Au travers de TD et de cours, j'ai
pu découvrir de nouvelles méthodologies d'apprentissage. Par ailleurs, j'ai pu présenter les
enseignements proposés au département GMP de l'IUT1 et plus généralement les possibilités
d'accueil d'élèves issus du département Mechanical Engineering du CIT sur le site Grenoblois
dans le cadre de stages.
16
Synthèse du Curriculum Vitae
1998
Postes et
Laboratoires
Thématiques
de recherche
1999
2000
Doctorant -Laboratoire de
Rhéologie - UBO
Suspensions d'argile
Polymère associatif
2001
2003
2002
Post-doctorat
ICE-HT
FORTH
ATER -Laboratoire
de Rhéologie - UBO
Polyampholytes
Polymère thermoassociatif
Emulsion
Mélanges
2004
2005
2006
2007
2008
2009
2011
2010
Maître de conférences Laboratoire
Rhéologie et Procédés
Post-doctorat
L2PIC
Thermoplastique
biodégradable
Renfort de membranes pour piles
Electrospinning
J. Rheol.
Publications
Enseignements
Responsabilités
pédagogiques
Langmuir
2 J. Rheol.
Polymer
SPE journal
Mécanique
Mécanique des fluides
IUT GMP Brest
Macromol.
Poly. Mat.
Sci. Eng.
PRL.
Macromol.
Soft Matter.
J. Rheol
Rheol Acta
Chap.3
Langmuir
Poly. Eng. Sci
Electrochim
Acta
Cellulose
Mécanique – Mathématiques - IUT GMP Grenoble
Rhéologie M1 GSI
Mécanique
GMP 1ère année
Projet GMP
1ère année
UE
GSI
Poursuites
d'études
GMP 2
Chapitre 2
Activités de recherche
Contribution à la rhéophysique et la mise en forme de polymères
1. Résumé des activités de recherche
Mon activité de recherche depuis ma thèse est orientée, de façon générale, vers la
rhéophysique expérimentale des fluides complexes. Ces fluides se structurent spontanément au
repos sur des distances caractéristiques supérieures à l'échelle de taille des constituants. Ces
structures complexes sont sensibles à l'écoulement et dépendent à la fois de la nature des
interactions entre constituants (électrostatique, stérique, Van der Waals, hydrophobe, liaison
hydrogène, …) et de leur intensité. Cette structuration modulable est à l'origine de
comportements mécaniques non linéaires originaux.
Mon travail de recherche a pour but de comprendre et de contrôler les propriétés
rhéophysiques de fluides complexes par une approche microstructurale. Je m'intéresse tout
particulièrement à la relation entre le comportement rhéologique macroscopique de fluides, leurs
microstructures et les interactions physico-chimiques mises en jeu à l'échelle des constituants.
Les méthodes de caractérisation rhéologique employées couvrent l'ensemble des essais
rhéométriques standards (écoulement permanent, oscillatoire, fluage, relaxation de contrainte).
Les résultats rhéologiques sont couplés à des analyses microstructurales à différentes échelles du
matériau par des moyens de microscopie optique, électronique à balayage, à force atomique,
mesures de diffusion dynamique et statique de la lumière, diffusion de neutrons aux petits angles.
Bien que mes différentes expériences professionnelles d'ATER, de post-doctorat et
maintenant de maître de conférences m'ont conduit à étudier des fluides aussi divers que des
suspensions d'argiles, des vases, des émulsions d'huile à base d'eau, des bruts lourds, j'ai
développé deux thématiques de recherche principales basées sur l'étude de deux classes de
matériaux polymères que je détaille dans ce manuscrit. Ces matériaux interviennent classiquement
dans la formulation de produits pharmaceutiques, cosmétiques, des peintures (polymères
associatifs) ou ils appartiennent au domaine des matériaux composites (thermoplastiques
chargés à base de matériaux biosourcés).
Première thématique de recherche : Rhéo-physique de polymères associatifs
[5], [7], [8], [9], [12], [15]
L'étude des polymères associatifs constitue un champ de recherches très important en
raison de l’étendue des applications de ces systèmes complexes. Ces matériaux polymères sont
principalement utilisés pour leur caractère épaississant dans l’élaboration de produits formulés
dans l’industrie des cosmétiques, dans l’industrie pharmaceutique, l'industrie des peintures ou en
papeterie par exemple. Plus généralement, ils sont utilisés comme agents de régulation des
propriétés rhéologiques de produits. Le caractère épaississant particulièrement prononcé de ces
polymères est lié à l'auto agrégation spontanée et réversible des chaînes de polymères en milieu
sélectif, résultant d'interactions attractives inter-chaînes entre groupes fonctionnels. Ces
interactions de faible énergie peuvent être brisées par un apport d'énergie de quelques k BT (kB, la
constante Boltzmann et T est la température absolue). Cette faible énergie d'interaction explique
le temps de vie court des jonctions associatives, variant de la microseconde à la milliseconde. Les
19
Résumé des activités de recherche
polymères associatifs peuvent former alors un réseau transitoire ou gel physique dont la signature
rhéologique est étroitement liée à la topologie du réseau.
La diversité de ces systèmes associatifs, en terme de composition chimique et d'architecture
des chaînes moléculaires, conduit à des structurations et donc des comportements rhéologiques
très variés. On peut toutefois classer ces matériaux en deux grandes catégories en fonction de
l'architecture des chaînes de polymère :
a) Les polymères hydrophiles neutres portant des groupes fonctionnels
associatifs. Ces polymères associatifs représentent la très grande majorité des systèmes étudiés
jusqu'à présent. Les groupes fonctionnels sont généralement des chaînons hydrophobes qui
peuvent être distribués le long des chaînes moléculaires ou localisés à leurs extrémités. Ces
derniers, appelés téléchéliques, ont fait l'objet de nombreuses études expérimentales1,2,
théoriques3,4 et de simulations .5,6 La structuration des chaînes en micelles de type fleur percolante,
introduite par Winnik et al7. (Fig. 1), serait responsable du comportement rhéo-épaississant
observé pour des contraintes intermédiaires (Fig. 2). Ce comportement est attribué :
- A la densification du réseau
- A l'étirement non-gaussien des chaînes
Fig. 1: Mécanisme d'auto-association de
polymères téléchéliques
Fig. 2: Comportement de solutions de polymères
téléchéliques sous cisaillement (Tam et al.8)
Pour les polymères associatifs en peigne, les groupes hydrophobes sont répartis le long de la
chaîne hydrophile. Ces polymères présentent une richesse de comportements rhéologiques
dépendant du taux de substitution de groupes hydrophobes, de leur distribution (blocs,
statistique, alterné, …), qui les distingue notablement des polymères téléchéliques. Ces deux types
de polymères (téléchélique et polymères associatif en peigne) représentent la majorité des
polymères associatifs étudiés jusqu'à présent.
T. Annable, R. Buscall, R. Ettelaie and D. Whittlestone, J. Rheol., 1993, 37, 695.
J.-F. Berret, Y. Sérero, B. Winkelman, D. Calvet, A. Collet and V. Viguier, J. Rheol., 2001, 45, 477.
3 F. Tanaka and S.F. Edwards, J. Non-Newtonian Fluids Mech., 1992, 43, 247, 273, 289.
4 A. Vaccaro and G. Marruci, J. Non-Newtonian Fluid Mech., 2000, 92, 261.
5 J. Huh, A.O. Balazs, J. Chem. Phys., 2000, 113, 2025.
6 T. Koga and F. Tanaka, Eur. Phys. J. E, 2005, 17, 115.
7 M. A. Winnik et al., Langmuir, 1993, 9, 881.
8 K.C. Tam K. C., Jenkins R. D., Winnik M. A., Bassett D. R., Macromolecules, 1998, 31, 4149.
1
2
Dans cette catégorie de polymères associatifs, je me suis intéressé à un nouveau polymère
thermoassociatif en peigne constitué d'une chaîne centrale de carboxyméthylcellulose (CMC), le
long de laquelle ont été greffés des groupes poly(N-isopropylacrylamide), (PNIPAM). Ces
groupes thermosensibles, synthétisés par voie radicalaire par l'équipe du Pr. Staikos du
département of Chemical Engineering de Patras, Grèce, ont la particularité de s'agréger
spontanément et de façon réversible au-dessus d'une température critique de 33°C. La formation
du réseau transitoire est basée sur la séparation de phase à l'échelle microscopique des chaînons
de PNIPAM au-dessus de cette température critique. J'ai caractérisé en particulier le
comportement viscoélastique en régime linéaire et non-linéaire de ce polymère à différentes
concentrations et températures.
Des mesures en écoulement permanent, en écoulement oscillatoire et en relaxation de
contrainte ont permis de mettre en évidence une transition de type gel faible/gel fort au-dessus
d'une température légèrement supérieur à la température d'agrégation des chaînons de PNIPAM.
Cette température de transition diminue avec l'augmentation de la concentration en polymère.
L'intérêt de ce type de polymère réside à la fois dans le caractère biocompatible de la chaîne de
CMC et des chaînons PNIPAM ainsi que le caractère associatif du PNIPAM observé sur une
large gamme de pH. Cette transition réversible au voisinage de la température corporelle fait de ce
polymère un excellent candidat pour la vectorisation de principes actifs. En effet, les microgels de
PNIPAM permettraient l'encapsulation et la libération prolongée de principes actifs dans des
environnements à pH variable (par voie orale ou sous forme de pansements).
b) Les polymères hydrophiles à caractère ionique portant des groupes
associatifs chargés ou neutres forment une classe de polymères associatifs plus
"confidentiels". Ces matériaux polymères ont des propriétés spécifiques dues à leur caractère
polyélectrolyte et/ou polyampholyte, que l’on retrouve dans un grand nombre de
macromolécules biologiques telles que les protéines. Ces matériaux sont très fortement stimulidépendants et offrent, de ce fait, un intérêt majeur sur le plan fondamental. En fonction de leur
constitution chimique, leurs propriétés rhéologiques dépendent fortement de paramètres
environnementaux comme le pH, la force ionique, la température, …
Dans cette catégorie de polymères, je me suis intéressé aux polyélectrolytes associatifs en
étudiant l'influence de la nature des interactions associatives.
- Considérons dans un premier temps le cas des polyélectrolytes téléchéliques qui se
distinguent très nettement des polymères téléchéliques classiques (chaîne hydrophile non ionique)
par une rigidité locale de la chaîne centrale sensible aux stimuli externes tels que le pH et la force
ionique. Pour cette étude, le copolymère à blocs est constitué d'une chaîne polyélectrolyte
principale de poly(dimethyl amino ethyl methacrylate), PDMAEMA, aux extrémités de laquelle
sont greffés des groupes poly(methyl methacrylate), PMMA hydrophobes. Ce polymère associatif
a été synthétisé par polymérisation "vivante" par l'équipe du Pr. Tsitsilianis du Department of
Chemical Engineering de Patras, Grèce.
Au-delà d'une concentration critique et pour un pH ~ 4, un réseau physique transitoire se
forme par association des groupes hydrophobes, conduisant à un comportement de type gel. Les
interactions coulombiennes contrôlant à la fois la rigidité moléculaire des chaînes et les
interactions inter-chaînes via le pH, sont responsables des propriétés rhéologiques spécifiques
très originales de ce matériau (seuil apparent d'écoulement, pseudo plateau Newtonien pour des
contraintes intermédiaires, deux dynamiques de relaxation très distinctes, …). Les caractérisations
viscoélastiques ont permis de mettre en évidence les mécanismes de relaxation aux temps courts
attribuées à la durée de vie des jonctions associatives et les mécanismes de relaxation aux temps
long correspondant à la reptation des chaînes, freinée par enchevêtrements électrostatiques.
20
21
Résumé des activités de recherche
- En conservant la même structure à blocs mais en substituant les groupes hydrophobes
par des groupes chargés négativement sur une chaîne centrale portant des charges positives,
nous obtenons un nouveau polymère associatif de type polyampholyte. La rigidité locale de
la chaîne est alors conservée mais ce type de polymère se distingue des polyélectrolytes
téléchéliques par la nature des interactions associatives de type électrostatique. Le copolymère à
blocs de cette étude, noté PAA135-P2VP628-PAA135 est constitué d'une longue chaîne centrale de
poly(2-vinyl pyridine), (P2VP), aux extrémités de laquelle sont greffés deux chaînons d'acide
polyacrylique (PAA).
A pH ~ 4, un réseau transitoire se forme par interactions électrostatiques entre les groupes
terminaux chargés négativement et la chaîne centrale portant des charges positives. Le réseau
ainsi formé présente des effets de structuration sous cisaillement très prononcés, attribués à la
densification des interactions électrostatiques. Contre toute attente, un comportement
thermoassociatif réversible a également été mis en évidence. L'étude rhéologique de ce
phénomène, couplée à des mesures de diffusion dynamique de la lumière ont permis d'attribuer
cet effet au gonflement des chaînes avec l'augmentation de la température sous l'effet de
l'amélioration de la solubilité du polymère. Cette étude montre pour la première fois qu'un
polymère peut présenter un comportement thermoassociatif sans pour autant comporter des
groupes fonctionnels à LCST (lower critical solution temperature), seule raison connue à ce jour
pour induire l'effet thermoassociatif.
L'effet épaississant de ces polymères n'est observé que sur une gamme étroite de pH
proche de 4. Dans le cadre du développement de vecteurs de principes actifs destinés à
l'administration de médicaments par voie orale, le mécanisme d'encapsulation doit être optimum
dans les conditions de pH très acide rencontrées dans le système digestif. L'utilisation de
polymères associatifs conventionnels doit donc être substituée au profit de matériaux procurant
cet effet épaississant sur une gamme de pH plus acide, proche de 2. Une alternative possible à
l'utilisation de polymères associatifs réside à utiliser des mélanges de polymères complexant
par liaisons hydrogènes. Ces systèmes sont à la marge des polymères dits "auto" associatifs
pour lesquels l'association s'effectue entre chaînes de même nature car dans ce cas le complexe
apparaît par interaction entre chaînes de nature différente. Dans ce cadre, j'ai étudié la formation
de complexes entre des chaînes de poly(acrylic acid-co-2-acrylamido-2-methylpropane sulfonic
acid)-graft-poly(N,N-dimethylacrylamide), (P(AA-co-AMPSA)-g-PDMAM) et des chaînes d'acide
polyacrylique PAA. En régime dilué, les mesures de diffusion dynamique de la lumière et de
diffusion de neutrons ont montré la formation de nanoparticules constituées d'un cœur insoluble
de complexe PAA/PDMAM, entouré d'une couronne anionique de P(AA-co-AMPSA). En
régime semi dilué, des mesures de viscosité et de viscoélasticité linéaire ont mis en évidence une
transition sol/gel à pH < 3.75. Cette transition est d'autant plus prononcée que la masse
moléculaire du PAA est élevée et que le taux de greffage en PDMAM est important.
J'ai mené cette recherche pluridisciplinaire dans le cadre de mon post d'ATER à l'université
de Bretagne Occidentale et au cours de mon post-doctorat à l'Institute of Chemical Engineering
and High Temperature Chemical Processes (ICE/TH-FORTH) de Patras en collaborations avec
et le Department of Chemical engineering de l'Université de Patras.
Deuxième thématique de recherche : Mise en forme et caractérisation de composites
[1], [2], [3], [4]
L'émergence des polymères composites est née de la nécessité d'optimiser les propriétés
d'usage des matières thermoplastiques (renfort mécanique, effet barrière, réduction de masse, …).
Les composites sont constitués d'une matrice polymère dans laquelle est dispersée une phase non
miscible. L'obtention de polymères composites peut se faire selon deux stratégies: par la
dispersion d'un ou plusieurs polymères ou par l'apport de charges solides dans la matrice.
Dans le cas de mélange de polymères, l'obtention de bonnes propriétés d'usage nécessite de
fragmenter très finement la phase dispersée pour former des inclusions de taille généralement
comprises entre 1 et 10 m et d'abaisser l'énergie interfaciale entre la matrice et la phase dispersée
par l'ajout de compatibilisant. Au cours de l'élaboration du mélange et de sa mise en forme, la
microstructure du composite dépend de l'équilibre entre le mécanisme de coalescence des
inclusions freiné par la présence du compatibilisant et le mécanisme de rupture des inclusions
induit par le taux de cisaillement8.
Dans le cas d'incorporation de charges solides, l'objectif principal est d'augmenter les
propriétés mécaniques des thermoplastiques à l'état solide. L'intérêt des charges solide et de
pouvoir atteindre des tailles de charge bien plus petites que dans le cas des mélanges de
thermoplastiques. Depuis l'émergence des nanomatériaux dans les années 1990, marquée par
l'arrivée des nanotubes de carbone, des nanofibres, la modification d'argile lamellaire, … de
nombreuses études ont été menées sur l'utilisation nanocharges comme renfort des
thermoplastiques. Elles ont permis de mettre en évidence le rôle crucial du rapport de forme des
charges9, de leur nature qui condition les interactions entre la charge et la matrice, leur
concentration ainsi que les conditions de mise en forme10.
Les polymères composites sont omniprésents dans les produits de la vie quotidienne et ils
sont devenus de ce fait indispensable. A l'heure actuelle, leur production repose à 99,5% sur
l'utilisation de polymères issus de la pétrochimie. La production de ces polymères représente à
l'échelle mondiale pas moins de 265 millions de tonnes en 2010 et adsorbe près de 4% de la
consommation de pétrole mondiale (source PlasticsEurope Market Research Group – PEMRG).
La mise en place récente d'initiatives environnementales visant à réduire les émissions de CO 2,
couplée à l'augmentation du prix du pétrole, a incité fortement l'utilisation de matériaux
biosourcés (matériaux issus de ressources renouvelables) dans la formulation des composites. Ma
thématique de recherche sur les composites s'inscrit pleinement dans cette démarche, avec deux
axes distincts:
- L'élaboration de bioplastiques à base d'amidon par extrusion réactive
-
Le renfort de membranes PEO par des charges cellulosiques
a) Elaboration de bioplastiques à base d'amidon par extrusion réactive
Cet axe de recherche a permis de démontrer la possibilité de mettre en œuvre un
bioplastique à base de poly(-caprolactone), PCL chargé en amidon aux propriétés rhéologiques
très proches de produits commerciaux naturels. Le caractère innovant réside ici dans l'élaboration
de ces matériaux en une passe par extrusion dite "réactive", combinant la dénaturation par
attaque acide des grains d'amidon et formation de formiate d'amidon, la compatibilisation du
formiate avec la matrice de PCL et le mélangeage. Les caractérisations des propriétés visqueuses
et viscoélastiques des mélanges, couplées à des observations en microscopie électronique à
balayage, ont permis d'optimiser les paramètres de formulation du mélange par l'amélioration
sensible de la compatibilisation de l'amidon dénaturé. L'influence du taux d'acide formique ainsi
que la nature et la masse moléculaire de l'oligomère utilisé comme plastifiant ont ainsi été mis en
évidence.
Huitric, J., Moan, M., Carreau, P.J. and Dufaure N., J. Non-Newtonian Fluid Mech., 2007, 145, 139.
Martone, A., Faiella, G., Antonucci, V., Giordano, M. and Zarrelli, M., Composites Science and Technology, 2011, 71, 1117.
10 Meier, J.G., Crespo, C., Pelegay, J.L., Castell,P., Sainz, R., Maser, W.K and Benito A.M., Polymer, 2011, 52, 1788.
8
9
22
23
Synthèse des principaux résultats
b) Renfort de membranes PEO par des charges cellulosiques
Cet axe de recherche s'inscrit dans le cadre du développement de batteries lithiumpolymère de très hautes performances nécessitant l'élaboration de nouvelles membranes
électrolytes polymères. L'objectif est d'optimiser les propriétés mécaniques de membranes à base
de poly(oxyde d'éthylène), PEO, de haute masse moléculaire, par ajout de charges naturelles tout
en conservant les propriétés de conduction ionique. Le choix des renforts naturels s'est porté sur
des celluloses de nature diverse, de type Coton, Ramie et Sisal et de morphologie différente avec
des microfibrilles, longues fibres flexibles et des whiskers, bâtonnets courts et rigides. L'influence
du procédé de mise en forme de ces membranes a également été étudié en comparent les
membranes obtenues par coulée/évaporation, procédé dit de laboratoire, et celles obtenues par
extrusion classiquement utilisée en industrie.
Les caractérisations rhéologiques et mécaniques des membranes obtenues par
coulée/évaporation ont mis en évidence un renfort mécanique très marqué pour de très faibles
taux de renfort. Dans le cas des membranes contenant des whiskers, ce renfort associé à la
formation d'un réseau percolant, est d'autant plus prononcé que le rapport de forme des
nanoparticules cellulosiques est élevé. En présence de microfibrilles, cet effet de renfort
mécanique, lié à la formation d'un réseau d'enchevêtrement des charges, est moins prononcé.
Pour les membranes extrudées, des observations en microscopie électronique à balayage
ont montré un effet de dégradation mécanique des charges et leur agrégation dans la matrice.
Contrairement aux films contenant des whiskers, une anisotropie des propriétés mécaniques des
films extrudés contenant des microfibrilles a démontré l'effet d'orientation des charges longues
induite par l'extrusion. Les propriétés rhéologiques de la matrice de PEO sont également
fortement affectées par la mise en forme et l'histoire mécanique du polymère. Un mécanisme de
scission des chaînes de PEO sous l'effet de l'écoulement élongationnel mis en évidence par des
mesures rhéométriques.
En présence de sel de Lithium, de type LiTFSI, le taux de cristallinité du PEO diminue, ce
qui affaiblit les propriétés mécaniques des films. Une légère diminution de la conductivité des
films renforcés et chargés en LiTFSI a été observée, très largement compensée par le
considérable renfort mécanique des membranes. Cette diminution des propriétés d'usage est liée
conjointement à la réduction de mobilité des chaînes de la matrice à l'interface PEO/cellulose, et
aux interactions entre le sel de lithium et le renfort cellulosique.
J'ai développé cette thématique au cours de mon second post-doctorat en collaboration
entre le Laboratoire de Rhéologie de l'Université de Bretagne Occidentale et le Laboratoire
Polymères, Propriétés aux Interfaces et Composites (L2PIC) de Lorient et plus récemment au
sein du Laboratoire Rhéologie et Procédés de Grenoble en tant que maître de conférences, en
collaboration avec le Laboratoire d'Electrochimie et de Physicochimie des Matériaux et des
Interfaces (LEPMI) et le Laboratoire Génie des Procédés Papetiers (LGP2). Cette dernière
collaboration m'a permis de co-encadrer la thèse d'Alessandra D'Aprea et le stage de master en
mécanique de Moulay Abdelkarim Elmoussaoui.
2. Synthèse des principaux résultats
2.1. Rhéo-physique de polymères associatifs
Les études que j'ai menées sur des polymères associatifs ont été guidées par une stratégie de
recherche peu courante, développée au Department of Chemical Engineering de l'Université de
Patras. Elle consiste à concevoir et synthétiser des structures moléculaires originales intégrant des
fonctions interactives, stimulables, autorisant le contrôle du processus associatif en vue de doter
les polymères synthétisés de propriétés rhéologiques spécifiques. Ces fonctions peuvent être
activées par une variation de la température (cas des polymères thermoassociatifs) ou par
modification du pH des solutions (cas des polyélectrolytes associatifs). Les variations de
température et de pH sont importantes dans le corps humain et ces types de polymères peuvent
être utiles dans la formulation de nouveaux médicaments.
2.1.1. Propriétés thermo-associatives de copolymères en peigne
2.1.1.1.
Introduction
Dans leur grande majorité, les polymères sont utilisés dans l'industrie comme additif
épaississant permettant de contrôler efficacement la rhéologie des produits. Ces polymères en
solution voient généralement leur viscosité décroitre selon une loi de type Arrhenius lorsque la
température augmente. Cette perte d'efficacité de l'effet épaississant en température conduit à une
perte du contrôle du procédé industriel dans beaucoup d'applications pour lesquels les fluides
sont soumis à des températures élevées (forage profond, couchage, production de produits
alimentaires, …). L'utilisation de polymères dont le caractère associatif est stimulé par
l'augmentation de température, appelés encore polymères thermoassociatifs, permet de pallier cet
inconvénient majeur.
Le concept de polymères thermoassociatifs, introduit par Hourdet et al.11, 12 au début des
années 1990, est basé sur les propriétés réversibles de polymères présentant une température
critique minimum de solubilisation (Lower Critical Solution Temperature ou LCST). Au-dessus
d'une température d'association Tassoc.= 32°C, ces polymères ne sont plus solubles et s'agrègent,
conduisant à une séparation de phase macroscopique. Les polymères thermoassociatifs sont
obtenus en introduisant dans, ou le long d'une chaîne fortement hydrophile, des groupes de
polymères à LCST. La séparation de phase se limite alors à des micro-domaines servant de points
de réticulation réversibles des chaînes hydrophiles.
Un grand nombre d'études a porté sur l'utilisation de poly(N-isopropylacrylamide),
(PNIPAM) comme groupe thermosensible. En effet, ce polymère biocompatible, non-toxique et
dont la température d'agrégation est proche de la température physiologique est utilisé dans le
domaine médical comme délivrance de médicaments pour le traitement de tumeurs solides,
couche de protection pour les médicaments, micelles pour une délivrance contrôlée de
médicaments et en ingénierie tissulaire pour produire des surfaces d’attachement/détachement de
cellules13. En revanche, peu de polymères d'origine naturelle ont été modifiés par greffage de
chaînes PNIPAM pour leur conférer un caractère thermoassociatif. L'étude porte ici sur les
propriétés visqueuses et viscoélastiques d'un CMC-g-PNIPAM constitué de
carboxyméthylcellulose, CMC, le long duquel sont greffés des chaînes latérales de PNIPAM. Le
choix du CMC est justifié par son utilisation commune comme excipient en galénique.
2.1.1.2.
Transition gel faible/gel fort des solutions de polymères thermoassociatifs
A la concentration de 6% en masse, les mesures de viscosité en cisaillement permanent,
représentées par les symboles pleins sur la figure 3, montrent très clairement une transition de
type fluide visqueux/fluide à seuil entre 37,5°C et 40°C. Cette transition se caractérisant par une
Hourdet, D., L'Alloret, F. and Audebert, R. Polym. Prepr, 1993, 34, 972.
Hourdet, D., L'Alloret, F. and Audebert, R. Polymer, 1994, 35, 2624.
13
Vihola, H., Laukkanen, A., Valtola, L., Tenhu, H. and Hirvonen, J., Biomaterials, 2004, 26, 3055.
11
12
24
25
divergence asymptotique de la viscosité au voisinage de la contrainte seuil. Ces mesures ont été
complétées par des mesures de fluage permettant d'accéder à la viscosité des solutions à très
faible taux de cisaillement en considérant la pente de l'évolution temporelle de la déformation du
fluide dans sa partie linéaire. Ces mesures montrent en fait qu'il s'agit d'une contrainte seuil
apparente, un plateau Newtonien étant observé aux faibles taux de cisaillement. Notons qu'un
comportement de type fluide à seuil pour des polymères en solution est rarement observé. Ce
comportement, également observé pour des HASE14, des chitosans15 et des gommes guar16
modifiés hydrophobiquement ainsi que pour des mélanges de HASE/tensio-actifs17, semble être
lié à la présence de micro-agrégats en suspension dans la solution de polymère.
Fig. 3: Viscosité d'une solution de CMC-g-PNIPAM à
C=6% en masse pour différentes températures.
Fig. 4: Viscosité Newtonienne pour des solutions de
CMC-g-PNIPAM à différentes concentrations en
fonction de la température.
La température de transition, notée T', est clairement mise en évidence en reportant les
valeurs des viscosités newtoniennes 0 en fonction de la température. Nous remarquons en effet
sur la figure 4 deux régimes de températures dont T' marque la limite:
-
Le régime de températures T < T', la viscosité newtonienne augmente exponentiellement
( ) . Ce régime de température correspond à un
avec la température selon la loi
régime de ségrégation faible pour lequel les interactions hydrophobes sont faibles,
conduisant à la formation d'agrégats lâches interconnectés. Dans ce régime de
température, une augmentation de concentration ou de température renforce très
efficacement le réseau transitoire.
Pour le régime de températures T > T', la loi de croissance de 0 avec la température est
( ) ( ) , avec  un paramètre qui décroit lorsque la concentration augmente.
Autrement-dit, l'augmentation de l'effet thermoassociatif marque une certaine saturation
d'autant plus marquée que la concentration augmente. Dans ce régime de températures
élevées, les interactions hydrophobes sont fortes et le renfort du réseau transitoire lié à
l'augmentation de concentration ou de température est amoindri.
Nous constatons également que la température de transition T' diminue significativement
lorsque la concentration augmente.
-
English, R.J., Raghavan, S.R., Jenkins, R.D. and Khan, S. A., J. Rheol, 1999, 43, 1175.
Esquenet, C. , Terech, P., Boué, B, and Buhler, E, Langmuir, 2004, 20, 3592.
16
Aubry, T. and Moan, M., J. Rheol. 1994, 38, 1681.
17
Tirtaatmadja, V., Tam, K. C. and Jenkins R. D., AIChE Journal, 1998, 44, 2756.
14
15
Cette transition apparaît également de façon très nette dans le comportement viscoélastique.
En cisaillement oscillatoire, la réponse rhéologique à un balayage en amplitude de déformation
est très classique dans le régime de
température T < T': au-delà de la zone
linéaire, les modules G' et G" diminuent.
En revanche, dans le régime de température
T > T', la zone linéaire se termine pour le
module G" par un pic (Fig. 5) alors que le
comportement du module G' reste inchangé
et décroit de façon continu avec l'amplitude
de déformation. L'amplitude du pic
augmente avec la température. Bien que
cette réponse se trouve hors du domaine
linéaire, les contributions des harmoniques
restent faibles et un tel pic doit être
considéré comme une réponse du matériau.
En effet, le rapport entre la contribution de
Fig. 5: Module de perte en fonction de l'amplitude
la réponse du signal fondamental à
ère
de déformation à C=6% en masse
l'harmonique d'ordre 3 (1 harmonique)
est inférieure à 1% au maximum du pic.
Comme nous le verrons dans la suite du manuscrit, un tel effet extra-dissipatif n'est pas rare pour
les polymères associatifs et n'est pas exclusif à ce type de matériau. J'ai pu en effet l'étudier en
détail dans le cas de suspensions concentrées de kaolin18. Bien qu'une interprétation commune
soit difficile, voire impossible à établir, des similitudes dans cette réponse viscoélastique
apparaissent entre ces systèmes. L'intensité de ce pic du module G" augmente lorsque l'intensité
des interactions entre constituants s'intensifie. Ceci s'observe lorsque la concentration augmente
pour les deux systèmes et lorsque la température augmente dans le cas du polymère associatif ou
lorsque la force ionique diminue (augmentation des répulsions électrostatiques inter particulaires)
pour la suspension de kaolin.
La signature du pic du module G" se retrouve dans les mesures de relaxation des contraintes
de la figure 6 à T = 45°C. Soumis à une déformation dans le régime faiblement linéaire du pic du
module G", la relaxation des contraintes
s'effectue en deux étapes: Dans une première
phase aux temps courts, les contraintes ne
relaxent pas. Il faut attendre un temps critique,
dont la valeur dépend de l'amplitude de
déformation imposée, pour qu'un mécanisme
de relaxation comparable à celui observé en
régime linéaire se mette en place. Notons que
ce temps critique évolue avec l'amplitude de
déformation de façon tout à fait similaire à
celle de l'intensité du pic du module G".
Comme l'a proposé Tirtaatmadja et al.14 Pour
des HASE, l'étirement des chaînes de
polymères pourrait à la fois expliquer le pic du
module G" l'absence de relaxation des
Fig. 6: Relaxation des contraintes en fonction du temps
contraintes au temps court dans le régime
pour une solution à 6% en masse à T = 45°C
faiblement non-linéaire.
18
Bossard, F., Moan, M. and Aubry, T., J. Rheol. 2007, 51, 1253.
26
27
2.1.2. Rhéo-physique de polyélectrolytes téléchéliques
2.1.2.1.
Introduction
Les polymères téléchéliques traditionnels, généralement constitués d'une chaîne centrale
de PEO aux extrémités de laquelle sont greffés des groupes hydrophobes de type alkyle ont fait
l'objet d'un nombre important d'études rhéologiques. Le schéma d'auto organisation de ces
polymères en boucle s'associant pour former des micelles de type fleur qui s'interconnectent pour
former un réseau tridimensionnel lorsque la concentration augmente est couramment admis. Je
me suis intéressé à une nouvelle classe de polymères téléchéliques pour lesquels la chaîne centrale
est un polyélectrolyte, en l'occurrence un poly(dimethyl amino ethyl methacrylate), PDMAEMA.
A pH acide, le caractère cationique du PDMAEMA confère à la chaîne une certaine rigidité qui
modifie la structure du réseau transitoire et sa dynamique moléculaire. En effet, comme le montre
les observations en microscopie à force atomique de la figure 7, les répulsions électrostatiques
interagissant le long de la chaîne hydrophile relativement courte (degré de polymérisation de 224
contre 32 pour les groupes hydrophobes) évitent la formation d'association intramoléculaire en
régime dilué, contrairement aux polymères téléchéliques standards.
Fig. 7: Représentation schématique du mécanisme d'association du polyélectrolyte
téléchélique issue d'observations AFM pour des concentrations croissantes.
De même, lorsque la concentration augmente, les chaînes ne s'associent pas en fleurs mais
forment des micelles en étoile qui s'interconnectent en régime concentré. Ces différences
majeures dans la structuration du réseau transitoire, associées à la présence d'interactions de
Coulomb entre chaînes chargées, confèrent au réseau ainsi formé une signature rhéologique
radicalement différente des polymères téléchéliques habituels. Enfin, la conformation de la chaîne
centrale dépend de son degré de neutralisation. L'intérêt de ce type de polymère réside donc dans
la possibilité de modifier la conformation de cette chaîne centrale en modifiant le pH et la force
ionique du milieu dispersant. L'étude rhéologique du réseau transitoire obtenu en régime
concentré a permis de mettre en évidence l'influence du pH et de proposer une interprétation à
l'échelle moléculaire des propriétés rhéologiques observées.
2.1.2.2.
Caractérisation des gels de polyélectrolytes téléchéliques à pH ~ 4
Les courbes d'écoulement des gels obtenus en régime concentré pour des pH ~ 4 ont des
comportements non-linéaires complexes, caractérisés par 4 zones d'écoulement (Fig. 8). La zone
I marquant la frontière entre deux zones de comportement Newtonien, en 0 et II, se traduit par
une chute brutale de viscosité de près de 4 décades. Cette discontinuité est, là encore, assimilable
à un comportement de type fluide à seuil apparent. Cette courbe d'écoulement se termine par un
comportement rhéofluidifiant classique aux fortes contraintes.
γ̇ max
Fig. 8: Viscosité et premières différences des contraintes normales d'une
solution à pH 3,5 et à la concentration de 1% en masse
La zone I est associée à la rupture des jonctions associatives. Nous avons observé une
fluctuation de la première différence des contraintes normales au cours du temps pour les 3
premières zones d'écoulement. Cette dépendance temporelle de N1 est également observée pour
les systèmes polymères à cristaux liquide et a été attribuée à l'oscillation des cristaux liquides
autour d'une position d'équilibre. Dans notre cas, ces fluctuations sont compatibles avec
l'oscillation de l'orientation moyenne des chaînes rigides autour d'une orientation privilégiée. La
zone III, marquant une stabilisation de N1 au temps long, correspond à l'alignement progressif
des chaînes dans la direction du champ d'écoulement.
Les mesures de viscoélasticité sous balayage en déformation montrent un pic du
module G" en fin de zone linéaire tout à fait similaire à celui observé dans le cas du polymère
thermo-associatif précédant (Fig. 9). Il est très intéressant de constater que la valeur du taux de
cisaillement correspondant au maximum du pic du module G", noté ̇
, de
-3 -1
l'ordre de 3.10 s , correspond aux taux de cisaillement marquant de début de la discontinuité de
viscosité de la Fig. 8. Outre ce comportement non-linéaire original, ces deux systèmes ont
également en commun un comportement de type fluide à seuil apparent.
Fig 9: Module viscoélastique d'une solution à 1% en masse
à pH 3,5 en fonction de l'amplitude de déformation
Fig 10: Relaxation de contraintes d'une solution
à C = 1% en masse pour différentes amplitudes
de déformation.
Les mesures de relaxation des contraintes de la Fig. 10 font apparaître deux modes de
relaxation bien distincts, très bien décrits par une fonction mono-exponentielle pour la relaxation
28
29
aux temps courts et une fonction exponentielle étirée pour le mode de relaxation aux temps
longs. Les deux temps de relaxation associés à ces modes sont respectivement de l'ordre de 0,1s
et 100s. Un tel mécanisme de relaxation avec deux modes a été observé pour des polymères
téléchéliques classiques mais seulement pour des amplitudes de déformation  > 200% bien audelà de la zone de linéarité. Dans le cas des polyélectrolytes téléchéliques, ce comportement
apparaît dès le régime faiblement non-linéaire ( > 0,7%). Ces deux familles de polymères
téléchéliques se distinguent également par la dépendance des deux temps de relaxation avec
l'amplitude de déformation. Lorsque l'amplitude de déformation augmente, les deux temps de
relaxation diminuent pour les polymères téléchéliques classiques, alors qu'ils évoluent de façon
opposée pour le polyélectrolyte téléchélique: le temps de relaxation court diminue contrairement
au temps de relaxation long qui augmente.
D'un point de vue moléculaire, le mode de relaxation aux temps courts a la même
origine pour les deux polymères téléchéliques, à savoir le désengagement de groupes
hydrophobes dans les jonctions associatives. Quant au mode de relaxation aux temps longs, il est
attribué au mécanisme de relaxation des chaînes freinées par les interactions électrostatiques des
chaînes voisines. Ces interactions électrostatiques sont de type répulsif entre les monomères mais
également de type dipôle attractif lié à la condensation des contres-ions sur les chaînes de
polyélectrolyte.
2.1.2.3.
Effet du pH sur les propriétés rhéologiques du polyélectrolyte téléchélique.
La figure 11 illustre l'influence du pH sur les propriétés rhéologiques du polyélectrolyte
téléchélique. La viscosité passe par un maximum au voisinage de pH 4. Dans cette gamme de pH,
le degré d'ionisation des chaînes de PDMAEMA est de l'ordre de 90%. Elles adoptent alors une
conformation étirée sous l'effet des répulsions électrostatiques ce qui peut produire deux effets
antagonistes sur le mode de relaxation court lié à la dynamique de désengagement des groupes
hydrophobes d'une jonction associative: La conformation étirée des chaînes contribue à diminuer
le temps de relaxation mais en contrepartie le temps de vie des jonctions associatif augment car il
est alors plus difficile pour un groupe hydrophobe désengagé d'une jonction associative de
s'impliquer dans une nouvelle jonction associative. Au-delà de pH 4, les chaînes se déprotonisent
et perdent graduellement leur rigidité alors que pour des pH < 4, les interactions électrostatiques
sont écrantées par l'augmentation de la force ionique des solutions.
Fig. 11: Viscosité newtonienne en fonction
du pH pour une solution à C = 1% en masse
Fig. 12: Module élastique (symboles pleins) et
modules de perte (symboles vides) d'une solution
à C = 1% en masse à pH 3,5 (carrés); pH 5,5
(ronds) et pH 7,5 (triangles)
Les mesures de viscoélasticité linéaire (Fig. 12) montrent que les deux modes de
relaxation sont dépendants du pH. Le mode de relaxation aux temps courts est mis en évidence
par l'augmentation des modules G" aux hautes fréquences et le mode de relaxation lent
correspond au maximum de la courbe Cole-Cole (G"=f(G')) dans sa partie semi-circulaire. Cette
réponse aux temps longs, proche d'un comportement de type Maxwell, est très polydisperse,
contrairement. Les deux temps de relaxation augmentent lorsque le pH diminue entre 7,5 et 5,5.
Cet effet est lié conjointement à:
- l'augmentation de la rigidité des chaînes qui augmente le temps de vie des
jonctions associatives.
- la condensation de contres-ions qui favorise les attractions entre chaînes et
freine la dynamique de relaxation au temps longs.
Cette étude montre l'intérêt que peut avoir ce type de polymères biocompatibles aux
propriétés dépendantes du pH dans le développement de médicaments. Sur le plan académique,
les résultats pourront être utilisés pour valider les modèles de réseaux transitoires des polymères
associatifs chargés.
2.1.3. Mécanismes d'association des polyampholytes
2.1.3.1.
Introduction
Les polyampholytes sont des polymères de la catégorie des polyélectrolytes, ayant pour
particularité de porter des groupes cationiques et anioniques au sein d'une même chaîne. A ce
titre, ils forment une catégorie particulière de matériaux de la classe des polymères associatifs. En
effet, alors que le caractère associatif des polymères associatifs est dû à l'auto-agrégation de
groupes fonctionnels, généralement de type hydrophobe, il est lié à l'interaction électrostatique
entre groupes de charges différentes dans le cas des polyampholytes.
Je me suis intéressé à un polymère à blocs, noté PAA135-P2VP628-PAA135, constitué
d'une longue chaîne centrale de poly(2-vinyl pyridine), (P2VP) aux extrémités de laquelle sont
greffés des chaînons d'acide polyacrylique (PAA). Les indices représentent le degré de
polymérisation de chaque bloc. Ce polymère présente un diagramme de phase complexe en
fonction du pH. A pH basique, les chaînes forment des nanoparticules constituées d'un cœur
hydrophobe de P2VP entouré de chaînes de PAA chargées négativement. La gamme de pH
compris entre 4 et 6,5 correspond au point isoélectrique du polymère, les nanoparticules
précipitent. Pour des pH inférieurs à 4, la chaîne centrale est protonée et le polymère se comporte
comme un polyampholyte.
120nm
Fig. 13: Viscosité spécifique en fonction du pH des solutions à la concentration de 3% en masse. Les images AFM
montrent la formation de particules à pH 7 et les chaînes en conformation semi-étirée à pH 2
30
31
Mon intérêt s'est porté sur le comportement du polymère à pH 3,5 pour lequel
l'intensité des interactions intermoléculaires est maximum, comme en témoigne la figure 13 avec
un maximum de la viscosité spécifique. A pH 3,5, les groupes PAA et P2VP sont partiellement
chargés et des interactions de type liaisons hydrogènes renforcent les interactions de Coulomb.
2.1.3.2.
Effet de structuration sous cisaillement des solutions de polyampholyte
Plusieurs comportements originaux ont été observés aussi bien en cisaillement permanent
qu'en cisaillement oscillatoire et suggérant un
effet de structuration des chaînes de
polymères induite par l'écoulement.
Trois régimes de concentrations sont
clairement identifiables en fonction du
comportement des solutions en cisaillement
permanent (figure 14). En régime dilué, les
solutions ont un comportement newtonien.
En régime semi-dilué, pour des contraintes
croissantes
(symboles
pleins),
le
comportement
est
successivement
newtonien, rhéo-épaississant puis rhéofluidifiant et pour des contraintes
décroissantes (symboles vides), la viscosité
augmente jusqu'à dépasser la valeur initiale
aux faibles contraintes. En régime concentré,
Fig. 14: Courbes d'écoulement à différentes concentrations
le comportement rhéo-fluidifiant est plus
marqué et la viscosité en fin de cycle de
contraintes est comparable aux valeurs initiales.
L'histoire mécanique imposée au système est primordiale dans le comportement des
solutions. En effet, l'augmentation de viscosité en régime semi-dilué n'apparaît que si le cycle de
contraints a dépassé le régime linéaire comme le montre la figure 15. Pour comprendre
l'augmentation de viscosité observée en fin de boucle de contrainte en régime semi-dilué, une
seconde boucle de contrainte, représentée sur la figure 16 par les symboles carrés, a été appliquée
après la 1ère boucle (symboles ronds). Pour des contraintes à nouveau croissantes, les solutions
n'ont plus de comportement rhéo-épaississant et la viscosité en fin de second cycle de contraintes
atteint une valeur comparable à celle en début de cycle.
Fig. 15: Viscosité en fonction de la contrainte pour une
solution à 4% en masse soumis à un cycle de contrainte
a) dans le régime linéaire et b) jusqu'au régime rhéoépaississant
Fig. 16: Viscosité en fonction de la contrainte pour
une solution à 4% en masse. Les symboles ronds et
carré représentent respectivement les résultats du 1er et
2ème cycle de contraintes.
32
En cisaillement oscillatoire, nous retrouvons à nouveau un pic du module G" en régime
faiblement non-linéaire mais ce pic se retrouve également pour le module G'. Pour ce dernier,
l'intensité du pic diminue avec l'augmentation de concentration. A notre connaissance, seuls les
polymères associatifs de type HASE présentent une telle dépendance des modules viscoélastiques
avec l'amplitude de déformation.
Fig. 17: Modules viscoélastiques des solutions à 4% et
6% en masse en fonction de l'amplitude de
déformation.
Fig. 18: Module élastique réduit en fonction de
l'amplitude de déformation aux concentrations de
3,5%(), 4%(), 4,5%(), 5%() et 5,5%() en
masse
L'interprétation des propriétés rhéologiques des solutions de polyampholytes à l'échelle
moléculaire repose sur une structuration particulière des chaînes de polymères au repos. Les
chaînes de polymères s'associent par interactions électrostatiques entre les groupes de PAA
chargés négativement et les chaînes centrales de P2VP chargées positivement pour former des
agrégats lâches en régime dilué, comme le montre le schéma de la figure 19 a.
En régime semi-dilué, un mécanisme de percolation des chaînes par associations
électrostatiques conduit probablement à former un réseau lâche de chaînes mécaniquement actifs
(Fig 19 b). Contrairement aux polyélectrolytes téléchéliques à la conformation de chaînes étirée, la
gamme d'amplitude de déformation du régime linéaire est ici très étendue (Fig. 17) suggérant que
les chaînes de polymères sont modérément étirées. Cette analyse est confirmée par les clichés
d'AFM de la figure 13. En effet, contrairement aux polymères téléchéliques, les bouts de chaîne
de PAA sont hydrosolubles et peuvent être stabilisées dans le milieu dispersant par répulsion
électrostatiques par les groupes de PAA voisins, formant de ce fait des branches pendantes et
mécaniquement inactives du réseau.
Fig. 19: Représentation schématique de l'organisation microstructurale des chaînes en régime a) dilué, b) semidilué et c) concentré.
33
En cisaillement oscillatoire, une amplitude de déformation croissante peut conjointement
favoriser ces branches pendantes à joindre le réseau mécanique et étirer les chaînes du réseau. Le
module élastique, dépendant à la fois de la densité de chaînes actives du réseau et de leur
élasticité, se trouve ainsi augmenté jusqu'à la rupture du réseau, ce qui peut expliquer le pic du
module G'. En cisaillement permanent, les contraintes de cisaillement correspondant au
comportement rhéo-épaississant ont le même effet structurant, incitant des interactions
intramoléculaires et les bouts de chaînes pendantes à former de nouveaux liens mécaniquement
actifs du réseau transitoire. Après déstructuration du réseau sous cisaillement élevé, les
associations intermoléculaires se reforment progressivement lorsque les contraintes de
cisaillement diminuent, aboutissant à la formation d'un nouveau réseau ayant une densité de
chaînes mécaniques plus élevée qu'à l'état initiale. La viscosité des solutions est alors plus élevée.
Soumis à un nouveau cycle de contraintes, seule s'établit la compétition entre l'association et le
désengagement des jonctions associatives. L'absence de densification du réseau sous cisaillement
peut alors expliquer l'absence de comportement rhéo-épaississant dans le second cycle.
En régime concentré, le confinement des chaînes réduit la densité de branches pendantes:
l'intensité du pic du module G' ainsi que l'intensité de l'effet rhéo-épaissiassant diminuent. Enfin,
la viscosité en début et fin de cycle de contraintes sont comparables.
2.1.3.3.
Etude du comportement thermo-épaississant
Nous avons découvert de façon fortuite que ce polymère présente un comportement
fortement thermo-épaississant et réversible. Comme l'illustre la figure 20, ce comportement se
manifeste entre autres par une variation de plus d'une décade des modules viscoélastiques entre
T = 10°C et T = 50°C. Un tel comportement est d'autant plus surprenant que ni la chaîne
centrale de P2VP, ni les chaînons de PAA ne présentent une température critique minimum de
solubilisation (LCST) comme les polymères thermo-associatifs classiques.
Fig. 20: Modules G' et G" en fonction de la température pour une solution de polymère soumise à un cycle
d'augmentation (symboles pleins) puis de diminution (symboles vides) de la température.
L'augmentation de température modifie également le comportement viscoélastique des
solutions. Lors d'un balayage en amplitude de déformation, l'amplitude du pic du module G'
augmente entre 20°C et 30°C puis diminue (figure 21). En revanche l'intensité du pic du module
G" ne fait qu'augmenter avec la température.
Fig. 21: Module élastique réduit G'/G'0 et Module de perte réduit G"/G"0 pour une solution à 4% en masse en
fonction de l'amplitude de déformation à T = 18 (), 22 (), 27.5 (), 30 (), 35 (), 45 (), and 50 °C ()
A la température de 12°C, le comportement viscoélastique linéaire en balayage en
fréquence des solutions à 4% en masse est typique d'un milieu dense de macromolécules. Il se
caractérise par une zone terminale aux basses fréquences et un croisement des modules
viscoélastiques aux fréquences élevées. Lorsque la température augmente, les modules
augmentent globalement et la réponse viscoélastique est qualitativement très proche de celle
observée en Fig. 12 pour le polyélectrolyte téléchélique.
Fig. 22: Modules élastiques (symboles vides) et modules
de perte (symboles pleins) d'une solution à 4 % en masse
en fonction de la fréquence à 12.5 (, ), 25 (,), et
50 °C (,).
Fig. 23: Temps caractéristique 𝜏𝑐
𝜋 𝜔𝑐 de la
solution à C = 4% en masse en fonction de la
température.
Le point de croisement des modules viscoélastiques apparaît à une fréquence c
associée à un temps caractéristique
. Ce temps court, de l'ordre de la seconde aux
basses températures, est compatible avec la durée de vie des jonctions associatives. Il augmente
brutalement au passage d'une température de 20°C pour atteindre plus de 100 secondes, puis
34
35
augmente de façon plus modérée jusqu'à 30°C, température au-delà de laquelle ce temps
caractéristique diminue. La discontinuité de la figure 23 au voisinage de T = 20°C marque la
transition sol/gel de la solution.
Les mesures de diffusion dynamique de la lumière, réalisées en régime dilué, ont permis
de déterminer le coefficient de diffusion D
des chaînes et d'en déduire leur rayon
hydrodynamique par la relation de StockesEinstein
(Fig. 24). Ces mesures
ont montré que le rayon hydrodynamique
des chaînes augmente entre T = 10°C et
30°C. Cette expansion des chaînes est
vraisemblablement due à une meilleure
solubilité des chaînons de PAA qui présente
une température critique supérieure de
solution
(Upper
Critical
Solution
Temperature ou UCST).
A faibles températures (T < 15°C)
les groupes de PAA sont proches d'une Fig. 24: Rayon hydrodynamique des chaînes de polymères
condition en solvant thêta et la possibilité à C = 0,5 % en masse en fonction de la température.
de former un réseau associatif est alors
faible. Lorsque la température augmente,
ces groupes s'étirent. Au voisinage de T = 20°C, un réseau de percolation se forme. Lorsque la
température augmente jusqu'à 30°C, les branches pendantes du réseau peuvent alors
progressivement intégrer le réseau de chaînes mécaniquement actives et ce réseau se rigidifie, ce
qui explique l'augmentation de l'intensité du pic module G'. Quant au module G", il reflète
généralement le volume occupé par le réseau. L'augmentation continue de l'intensité du pic du
module G" avec la température est en accord avec l'expansion continue des chaînes. Cette
expansion des chaînes est en compétition avec l'agitation thermique croissante. Au-delà de 30°C
les effets de l'agitation thermique prédominent: le coefficient de diffusion des chaines augmente
ce qui se traduit par une diminution artificielle de RH, une décroissance de la viscosité des
solutions selon la loi d'Arrhenius, et une fragilisation du réseau élastique qui explique la
diminution de l'amplitude du pic du module G'.
L'originalité de la réponse thermique du polymère réside dans l'absence de groupe à
LCST, qui était jusqu'à présent la seule origine connue du comportement thermo épaississant de
certaines solutions de polymères associatifs. L'expansion des chaînes, induite ici par
l'augmentation de température, est un paramètre clés souvent négligé dans l'étude des polymères
associatifs.
2.1.4. Rhéologie de mélanges de polymères complexant par liaison hydrogène
2.1.4.1.
Introduction
Comme nous l'avons vu précédemment, le caractère épaississant des polymères
associatifs provient généralement de l'interaction entre groupements hydrophobes de chaînes
voisines ou d'interactions électrostatiques de type Coulomb entre charges opposées. Les liaisons
hydrogènes, ou interactions d'origine électrostatique entre dipôles, peuvent également être à
l'origine du caractère épaississant de certains polymères. C'est le cas des mélanges de polymères
complémentaires comportant un polymère donneur de protons (acides carboxyliques faibles par
exemple) et un polymère accepteur de protons (poly bases non ioniques). Cependant, leur
utilisation est fortement limitée par la gamme réduite de pH dans laquelle ils sont solubles. En
effet, à pH > 4 - 5, l'augmentation de la densité de sites ionisés chez le donneur de protons
provoque une diminution de densité des liaisons hydrogènes: les complexes ne se forment pas.
En revanche, pour des pH inférieurs à 3 - 3,5, la fraction d'anions carboxylate responsable de la
solubilité du complexe diminue, conduisant à une précipitation du complexe.
Une extension de la solubilité de tels complexes jusqu'à pH = 2 ouvre des perspectives
d'applications nouvelles pour lesquelles un effet épaississant est requis à pH très acide. Une telle
extension de la solubilité du complexe de
polymère a été obtenue en synthétisant un
copolymère anionique greffé P(AA-coAMPSA)-g-PDMAM (Fig. 25), mélangé avec un
PAA dans un rapport 1:1. La présence de
groupes anioniques AMPSA fortement chargés
dans le copolymère greffé permet d'augmenter
de façon significative la solubilité de la chaîne
principale et d'empêcher la précipitation du
complexe à pH inférieur à 3,5. Les chaînes
pendantes de PDMAM hydrosolubles ont
d'importantes propriétés de donneur de protons Fig. 25: Structure du copolymère P(AA-co-AMPSA)et s'associent avec les chaînes de PAA par g-PDMAM
liaison hydrogène.
2.1.4.2.
Transition sol/gel en régime semi-dilué
A la concentration totale de 6% en masse (régime semi-dilué), l'étude a porté sur
l'influence du pH, du taux de greffage en PDMAM en % de la masse de P(AA-co-AMPSA) et de
la masse moléculaire du PAA sur les propriétés rhéologiques des mélanges P(AA-co-AMPSA)-gPDMAM et PAA.
L'étude de l'effet du pH a porté sur le mélange du P(AA-co-AMPSA)-g-PDMAM avec un
taux de PDMAM de 60% et un PAA de masse moléculaire
g/mol. A pH = 3,8, la figure
26 montre que le mélange se comporte comme un fluide viscoélastique avec une zone terminale
qui s'étend sur une gamme de fréquences élevées. Lorsque le pH diminue, les liaisons hydrogènes
entre PAA et PDMAM sont progressivement renforcées ce qui augmente les modules
viscoélastiques. A pH = 3,4, le point de croisement des modules viscoélastiques est atteint à la
pulsation de 10 rad/s et il se trouve décalé vers les faibles pulsations à pH = 2. Ce décalage
progressif du point de croisement des modules vers les faibles pulsations résulte du
ralentissement de la dynamique moléculaire induite par le complexe PAA/PDMAM. A pH = 2, le
mélange se comporte alors comme un gel avec G' > G" et les deux modules faiblement
dépendants de la pulsation. Un réseau transitoire est alors formé, combinant un complexe
insoluble PAA/PDMAM obtenu par liaison hydrogène et agissant comme point de contact entre
les chaines principales P(AA-co-AMPSA) anioniques et hydrosolubles.
36
37
Fig. 26: Modules viscoélastiques en fonction de la pulsation pour des mélanges 1:1 de P(AA-co-AMPSA)-g-PDMAM
et PAA à C = 6% à pH 3,8; 3,4 et 2 et un schéma des structure correspondantes de part et d'autre de pH = 3,75
ding structure
Pour étudier l'influence du taux de greffage en PDMAM, le pH est fixé à 2 avec le PPA à
la masse moléculaire de
g/mol. Trois copolymères anioniques avec des taux de greffage
en PDMAM de 22%, 42% et 60% en masse ont été synthétisés. L'augmentation du taux de
greffage provoque l'augmentation de la viscosité des mélanges et un renforcement de l'effet rhéofluidifiant. Pour le taux de greffage maximum de 60%, le profil des courbes d'écoulement est
comparable à celui observé en Fig. 14 pour les polyampholytes. Il se caractérise par un
comportement rhéo-épaississant pour des contraintes intermédiaires, une chute de viscosité de 4
décades au-delà d'une contrainte critique et une boucle d'hystérésis lorsque la contrainte diminue.
Comme dans le cas des polyampholytes à blocs, les chaînes de P(AA-co-AMPSA) ont une
configuration étirée et sont en répulsion mutuelle liée à la présence des groupes AMPSA. Cette
similitude suggère une structuration comparable du réseau de chaînes de polymères.
Fig. 27: Viscosité en fonction de la contrainte pour
des mélanges 1:1 de P(AA-co-AMPSA)-g-PDMAM et
PAA à C = 6% à pH = 2 pour des taux de PDMAM
de 22% (,), 42% (,) et 50% (, ). Les
symboles
pleins
et
vides
correspondent
respectivement à l'application de contraintes
croissantes et décroissantes.
Fig. 28: Viscosité en fonction de la contrainte
pour des mélanges 1:1 de P(AA-co-AMPSA)-gPDMAM/ PAA90 (,) et P(AA-co-AMPSA)-gPDMAM/PAA450 (, ) à C = 6% à pH = 2.
Les symboles pleins et vides correspondent
respectivement à l'application de contraintes
croissantes et décroissantes.
L'influence de la masse moléculaire des PAA a été mise en évidence à pH = 2, pour un
taux de greffage de 42% en masse. Deux PAA de masse moléculaire
et
g/mol ont
été utilisés. L'augmentation de viscosité avec l'augmentation de la masse moléculaire du PAA,
observé en figure 28, reflète la possibilité pour les chaînes de PAA plus longues de connecter plus
de chaînons de PDMAM.
A notre connaissance, ce système est le premier exemple de polymères formant un réseau
transitoire via uniquement des liaisons hydrogènes dont les propriétés rhéologiques sont
contrôlées par le pH et les paramètres moléculaires des différents polymères.
2.1.4.3.
Formation de nanoparticules en régime dilué
En régime dilué, nous nous sommes intéressé au complexe formé par le mélange du
copolymère P(AA-co-AMPSA)-g-PDMAM avec un taux de greffage en PDMAM est de 48% et
du PAA de petite masse moléculaire (
g/mol) avec des rapports molaires [PAA90] /
[PDMAM] compris entre 0,25 et 1,5. Dans ces conditions de concentrations et de paramètres
moléculaires, les complexes forment des nanoparticules dont la structure a été étudiée par
diffusion de neutrons aux petits angles, diffusion statique et dynamique de la lumière et
microscopie à force atomique.
Les mesures de diffusion de neutrons aux petits angles dans le régime de Guinier
qRg << 1, suggèrent la formation d'agrégats de taille finie (Fig. 29). Pour des vecteurs d'onde
intermédiaires, 0,01 < q < 0,1 Å-1, l'intensité diffusée diminue suivant une loi de la forme I ~ q- d.
La valeur de l'exposant d, comprise entre 3,5 et 4, suggère la présence d'objets tridimensionnels à
la surface fractale attribués aux complexes insolubles PDMAM/PAA. Aux valeurs de q > 0.1 Å-1,
la diffusion est vraisemblablement liée aux chaînes anioniques hydrosolubles, constituant une
coque autour du cœur. Une vue schématique du complexe est proposée en figure 29. L'intensité
diffusée est maximale pour un rapport r = 1,1, correspondant aux conditions stœchiométriques
des mélanges [PAA90] / [PDMAM].
Fig. 29: Intensité des neutrons diffusés en fonction du vecteur d'onde q pour un mélange à la concentration de
6,3.10-3g/cm3. Le schéma de droite représente l'organisation des chaînes de polymères dans la l'agrégat.
38
39
La décroissance de l'intensité des neutrons diffusés dans le régime de Guinier dépend du
rayon Rc du cœur hydrophobe selon la loi exponentielle suivante:
(
)
L'ajustement des courbes expérimentales montre que le rayon du cœur est de l'ordre de
17 nm. Les observations par microscopie à force atomique de la Fig. 30 confirment la présence
d'agrégats compacts de chaînes. En corrigeant de la taille de la pointe du cantilever, l'AFM
permet d'observer le cœur hydrophobe dont le rayon, de l'ordre de 22 nm, est en très bon accord
avec les mesures de diffusion de neutrons.
Fig. 30: Observation AFM d'un mélange P(AA-co-AMPSA)-g-PDMAM
/PAA déposé sur un support de mica.
Ces mesures ont été complétées par de la diffusion dynamique de la lumière, permettant
d'estimer le rayon hydrodynamique de la particule en prenant en compte la coque des chaînes
P(AA-co-AMPSA). L'extrapolation à concentration nulle du coefficient de diffusion a donné une
valeur D0 correspondant à un rayon hydrodynamique de 105 nm pour une particule isolée.
De telles mélanges de polymères interagissant par liaisons hydrogènes et formant des
nano-particules aux pH basiques sont des systèmes prometteurs pour modéliser des principes
actifs encapsulés, utilisés comme vecteurs pharmaceutiques.
2.1.5. Conclusion de l'étude de polymères associatifs
Les polymères associatifs constituent une catégorie de polymères hydrosolubles aux
comportements rhéologiques riches et variés. Outre les polymères associatifs hydrophobes
traditionnels, l'étude menée ici a montré l'intérêt des interactions électrostatiques (interaction de
Coulomb ou liaison hydrogène) intervenant à la fois comme jonctions associatives et contrôlant
la conformation des chaînes de polymères. Ces nouveaux polymères associatifs présentent une
réponse rhéologique modulable par des paramètres habituels que sont la concentration ou la
température mais également par la force ionique et le pH. Ces derniers paramètres permettent de
modifier simultanément l'intensité des interactions associatives et la rigidité des chaînes.
Lorsque les conditions d'interactions associatives sont optimales, le comportement
rhéologique de ces polymères se caractérise par une discontinuité dans les courbes d'écoulement,
traduisant un comportement de type fluide à seuil apparent. Ce comportement s'accompagne
également d'un pic du module G" dans les courbes de balayage en amplitude de déformation. Le
taux de cisaillement correspondant au maximum du pic du module G" semble coïncider avec
celui associé au début de la discontinuité de la courbe d'écoulement. Une question reste posée: le
pic du module G" est-il une marque du comportement de fluides à seuil apparent? Des
réorganisations locales du réseau, responsables de l'effet extra-dissipatif, pourraient avoir lieu
juste avant sa rupture.
2.2. Mise en forme et caractérisation de composites
Ma seconde activité de recherche, initiée depuis mon post-doctorat au L2PIC, concerne le
développement de composites biosourcés. Cette activité est centrée sur l'utilisation de charges
d'origine naturelle comme renfort de matrices polymères. Mon intérêt s'est porté sur l'influence
du procédé de mise en forme et sur les paramètres de formulation des composites (nature,
concentration des constituants) sur leurs propriétés d'usage. Je me suis intéressé tout
particulièrement à l'utilisation d'amidon dénaturé comme renfort de matrice de polycaprolactone,
PCL dans l'objectif de proposer un matériau de substitution aux matériaux de type Mater-Bi. J'ai
également étudié l'utilisation de fibres cellulosiques comme renfort de membranes de poly(oxyde
d'éthylène), PEO, dédiées aux piles Lithium.
2.2.1.
Elaboration de bioplastiques à base d'amidon par extrusion réactive
2.2.1.1.
Introduction
Les matériaux thermoplastiques d'origine fossile offrent de nombreuses propriétés (résistance
mécanique, déformabilité, faible densité, matériau hydrofuge, …) qui en font des matériaux
incontournables dans notre vie quotidienne. Nous les trouvons abondamment dans le domaine
de l'emballage, le textile, le bâtiment, le transport, des équipements électriques et électroniques,
… Néanmoins, avec une durée de vie de l'ordre de 200 ans pour la plupart des thermoplastiques
dérivés du pétrole, leur utilisation intensive pour des usages courants est à l'origine d'une
pollution des sols, des cours d'eau et des océans.
Depuis le début des années 1970, des travaux ont été menés pour développer des matériaux
polymères combinant les caractéristiques techniques des thermoplastiques de la pétrochimie et le
caractère biodégradable. Ceci peut être obtenu à partir de polymères biodégradables de synthèse
de type aliphatiques tels que le polycaprolactone, PCL. L'incorporation de matériaux d'origine
naturelle tels que l'amidon, bon marché (0,5 à 1 €/kg) et abondant (production annuelle
européenne de l'ordre de 6 millions de tonnes), permet de réduire le coût de production de ces
matériaux. Leur utilisation doit cependant être raisonnée, les terres agricoles devant servir en
priorité les besoins alimentaires. Cependant, l'affinité très élevée de l'amidon pour l'eau rend très
difficile l'élaboration de matériaux composites à base d'amidon. Il est possible de limiter
l'hydrophilie de l'amidon et d'augmenter son caractère thermoplastique tout en préservant sa
biodégradabilité en substituant ses groupements hydroxyles par des groupements ayant moins
d'affinité pour l'eau. Récemment, une telle modification chimique de l'amidon natif par
formiatation ou modification chimique de l'amidon en formiate d'amidon par attaque à l'acide
formique a été brevetée par le L2PIC. Cependant, l'élaboration d'un composite en deux étapes
(dénaturation de l'amidon en formiate d'amidon par voie liquide et extrusion du composite
PCL/formiate d'amidon) n'est pas viable à l'échelle industrielle. Dans le cadre du programme de
recherche inter régional
Amidon
AMIDODER,
mon
Compatibilisant
projet de recherche a
PCL
consisté à étudier la
Extrusion
modification chimique en
masse de l'amidon et sa
compatibilisation avec la
Acide formique
Dégazage
matrice de PCL par un
Eau
procédé
d'extrusion
réactive décrit en Fig 31.
Fig. 31: Déroulement schématique du procédé d'extrusion réactif
40
41
En optimisant le procédé d'extrusion, l'objectif était de proposer un nouveau composite aux
propriétés mécaniques proches des produits actuellement commercialisés, tel que la gamme
Mater-Bi proposée par la société Novamont, et pouvant être produits en grande quantité par des
moyens industriels classiques et à un prix compétitif.
Tous les mélanges réalisés contiennent 40% en masse d'amidon, 30% en masse de PCL et
40% en masse de compatibilisant de type oligomère. L'étude des composites à base de formiate
d'amidon a porté sur l'influence du rapport acide formique/amidon, la masse moléculaire du
compatibilisant, un 1,6-hexane-dioladipate et phtalates et sa nature chimique en substituant
l'oligomère précédant par un PCL de petite masse portant des groupes hydroxyles.
2.2.1.2.
Caractérisation rhéologique et microstructurale des composites
Etude d'un matériau commercial: Le matériau Mater-Bi ZF03UA, commercialisé par la
société Novamont et largement utilisé dans l'industrie des emballages, a été choisi comme produit
de référence pour ce projet. Ce matériau est principalement constitué d'une matrice PCL et
d'amidon dénaturé. Le cliché en microscopie électronique à balayage en la Fig. 32 d'une surface
de cryofracture montrent une phase co-continue du composite sans présence visible d'inclusion
d'amidon à l'échelle microscopique. La courbe d'écoulement de la Fig. 33, obtenue en
superposant des mesures de fluage pour les bas gradients, de cisaillement permanent (symboles
vides) et des mesures en cisaillement oscillatoire à 95°C pour les gradients plus élevés, est bien
décrite par la somme de deux modèles de Cross:
( ̇)
(
̇)
(
̇)
La contribution ( ̇ ) aux faibles taux de cisaillement est associée au plateau de viscosité 01,
de l'ordre de 108 Pa.s, et à une chute de viscosité dont la dépendance avec le taux de cisaillement
est proche de ̇ . Ce comportement aux faibles taux de cisaillement, témoignant d'un
comportement de type fluide à seuil apparent, peut être associé à la contribution visqueuse d'un
réseau percolé d'amidon finement déstructuré.
La contribution ( ̇ ), observée pour des taux de cisaillement élevés, se caractérise par un
plateau de viscosité 02 de l'ordre de 5. 104 Pa.s et correspond à la réponse rhéologique de la
matrice de PCL.
Fig. 32: Observation MEB d'une cryofracture
de ZF03UA
Fig. 33: Viscosité en cisaillement permanent (symboles vides)
et viscosité complexe (symboles pleins) en fonction du taux de
cisaillement et de la pulsation.
42
Etude des mélanges formiate d'amidon/PCL: L'étude de l'influence du rapport acide
formique/amidon a été réalisée en utilisant le 1,6-hexane-dioladipate et phtalates de masse
2700g/mol comme plastifiant. Les clichés MEB de la figure 34 montrent les surfaces de
cryofracture des composites (a) sans acide formique et avec un rapport acide formique/amidon
de (b) 30% et (c) 60%. Pour tous ces composites, l'amidon se trouve sous la forme de nodules de
taille microscopique. L'énergie mécanique de 300 kJ/kg imposée pendant l'extrusion n'est pas
suffisante pour déstructurer l'amidon comme c'est le cas pour le composite du commerce. L'acide
formique ne modifie pas sensiblement la taille des nodules d'amidon mais semble augmenter
l'affinité des grains d'amidon pour la matrice et simultanément attaquer chimiquement la matrice
de PCL.
(a)
(b)
(c)
Fig. 34: Clichés MEB de cryofractures de composites PCL/amidon (a) sans acide formique,
(b) avec 30% et (c) 60% d'acide formique
Cette observation se confirme sur les courbes d'écoulement de la figure 35, proches
qualitativement de celle du matériau de référence. Le plateau de viscosité aux faibles taux de
cisaillement 01 passe par un maximum pour un taux d'acide formique de 15%. En revanche, le
second plateau de viscosité, aux taux de
cisaillement élevés, décroît lorsque le taux
d'acide formique augment. Cet effet
s'explique par l'action combiné de l'acide:
-
sur
l'amidon,
qui
déstructure
partiellement la surface des grains
d'amidon et renforce de ce fait les
interactions entre les groupes ester et
hydroxyle de l'oligomère et les groupes
hydroxyle et formiate de l'amidon
modifié,
-
sur la matrice; les chaînes de PCL étant
sensibles à l'attaque chimique de l'acide.
Fig. 34: Viscosité des composites sans acide formique (B1-0) et
avec 15% (B1-15), 30% (B1-30), et 60% (B1-60) d'acide
formique.
Pour ce composite, le taux de 15% en acide formique semble être le taux optimal entre
une augmentation de la compatibilisation et l'affaiblissement de la matrice.
43
Pour les clichés de la figure 35 (a) et (b) des composites contenant respectivement 30% et
60% d'acide formique, l'oligomère utilisé est de plus grand masse, passant de 2700g/mol à
7400g/mol. La déstructuration de l'amidon et sa compatibilisation avec la matrice est alors plus
notable, aboutissant à la formation d'une phase pratiquement co-continue pour un taux d'acide de
60%. Cette augmentation de la compatibilisation se traduit sur la figure 36 par un maximum des
deux plateaux newtoniens, la contribution 01 des nodules et 02 de la matrice, pour un taux
d'acide formique de 30%. En effet, l'oligomère de plus grande masse se localise:
-
à l'interface format/PCL et favorise les interactions entre nodules
-
dans la matrice de PCL augmentant ainsi sa viscosité.
(a)
(b)
Fig. 35: Clichés MEB de cryofractures
de composites PCL/amidon avec
l'oligomère de forte masse pour (a)
30% et (b) 60% d'acide formique
Fig. 36: Viscosité des composites sans acide formique (B2-0) et avec
15% (B2-15), 30% (B2-30), et 60% (B2-60) d'acide formique.
Comme nous pouvons le constater, le compatibilisant joue un rôle déterminant dans la
morphologie et le comportement rhéologique du composite. L'effet compatibilisant a pu être
optimisé en choisissant un oligomère de PCL portant des groupes hydroxyle. Le renforcement de
la compatibilisation est dû aux interactions par liaison hydrogène entre l'oligomère modifié et le
formiate. En présence de 15% d'acide, la figure 37 montre que l'amidon est profondément
dénaturé; aucun grain n'est perceptible à l'échelle de quelques microns. Aux faibles taux de
cisaillement, la courbe de viscosité du mélange atteint alors un maximum proche de celle mesurée
pour le matériau de référence.
Cette étude a démontré la possibilité de produire un composite à base d'amidon à l'échelle
pilote, aux propriétés rhéologiques proches du matériau commercial de référence à l'état fondu.
La réaction de formiatation partielle de l'amidon a été réalisée en cours d'extrusion. Les sociétés
Europlastiques et Linpac ont proposé de continuer des essais à partir de ces nouveaux
matériaux.
Fig. 37: Clichés MEB de cryofractures
de composites PCL/amidon avec
l'oligomère de PCL portant des
groupes hydroxyles et avec 15%
d'acide formique
2.2.2.
Fig. 38: Viscosité des composites avec l'oligomère de type PCL
modifié, sans acide formique (B3-0) et avec 15% (B3-15) et 60%
(B2-60) d'acide formique.
Renfort de membranes PEO par des charges cellulosiques
2.2.2.1.
Introduction
Dans le cadre de la Thèse d'Alessendra D'Aprea, en collaboration avec le LEPMI,
(Laboratoire d’Électrochimie et de Physico-chimie des Matériaux et des Interfaces) et le LGP2
(Laboratoire Génie des Procédés Papetiers) nous avons étudié l'effet de renfort de fibres de
cellulose au service d'une technologie en développement: les batteries Lithium - polymère.
Le stockage d'énergie destiné aux applications nomades (téléphonie mobile, ordinateurs
portables, jeux, …) utilise très largement la technologie Lithium-ion. Un sel de Lithium dissout
dans un solvant organique permet le transfert des ions Li+ entre l'anode en graphite et la cathode,
un oxyde métallique. Ces batteries présentent une bonne conductivité, de l'ordre de 1 mS/cm à la
température ambiante, et permettent un fonctionnement entre -20°C et 60°C mais elles sont
instables dans certaines conditions d'utilisation. Elles peuvent effectivement s’enflammer
facilement en cas de choc, de surcharge électrique ou lors d’un assemblage défaillant. Ainsi, Dell,
Sony en 2005 et HP en 2011 ont rappelé un nombre important de batteries défectueuses. La
technologie dite Lithium-polymères, Li-Po, est peut-être en voie de supplanter les batteries à
lithium liquide. En effet, dans le cas des batteries Lithium-polymère, l'électrolyte liquide est
remplacé par une membrane polymère imbibée de sel de Lithium. Ceci offre deux avantages
principaux:

L'utilisation de boitiers rigides et étanches pour contenir l'électrolyte n'est plus
nécessaire. Les batteries Li-Po peuvent être fabriquées avec des enveloppes
plastiques plus légères, de formes plus complexes et plus fines permettant une
meilleure miniaturisation de la batterie. Ceci explique qu'elle équipe actuellement
les oreillettes Bluetooth par exemple.
44
45

Les membranes polymères ont moins de composants volatils et inflammables que
les électrolytes liquides. Enfin, le déplacement d'éventuelles impuretés métalliques
responsables de court-circuit est évité. Ces batteries sont donc plus sûres.
En contrepartie, les batteries Li-Po ont une faible densité énergétique. Une voie de
développement de cette technologie passe par l'optimisation de l'électrolyte polymère. Les
polyéthers possédant une bonne stabilité électrochimique, de bonnes propriétés de conduction
ionique et une bonne tenue mécanique sont des matériaux de prédilection pour cette application.
Le poly(oxyde d'éthylène), PEO, de haute masse molaire est un polymère très utilisé comme
séparateur du fait de sa capacité de solvatation du cation lithium. Cependant, la température de
fonctionnement des batteries est proche de la température de fusion du PEO (vers 60°C) et la
membrane perd ses propriétés mécaniques. Pour remédier à ce problème, nous avons considéré
l'utilisation de charges cellulosiques comme renfort mécanique des membranes. Le choix de ce
type de charges se justifie par leur faible densité, leur caractère renouvelable et leur disponibilité à
travers le monde sous formes variés, en quantité abondante et à des coûts de production faibles.
Enfin, l'utilisation de ces fibres permet aux composites d'être recyclés, contrairement aux
composites à base de fibres de verre, de kevlar et plus récemment de nanotubes de carbone.
Les objectifs de ces travaux visaient à formuler des électrolytes polymères
nanocomposites innovants présentant une bonne tenue thermomécanique et des bonnes
propriétés de conduction. Pour cela, la morphologie des charges (charges courtes ou whiskers,
charges longues ou microfibrilles), leur nature (sisal, coton, ramie) et le procédé de mise en forme
(coulée/évaporation ou extrusion) ont été considérés.
2.2.2.2.
Influence de l'histoire mécanique sur la matrice de PEO
Le choix de la matrice s'est porté sur un PEO de forte masse moléculaire (5.10 6g/mol),
permettant d'assurer des propriétés mécaniques élevées en l'absence de charges. Deux procédés
de mise en forme ont été utilisés: un procédé par coulée/évaporation permettant de fabriquer les
membranes de façon contrôlée à l'échelle laboratoire et le procédé d'extrusion, privilégié pour
une production à l'échelle industrielle.
Dans le cas des membranes obtenues par coulée/évaporation, l'étape préliminaire
consiste à solubiliser le polymère dans un solvant, ici l'eau distillée. Or, à concentration égale,
l'étude bibliographique montre des disparités dans la viscosité des solutions en fonction du mode
de solubilisation du polymère; disparités que nous avons constatées en comparant la rhéologie de
solutions dispersées par agitation au barreau magnétique (stirred solutions) ou secouées par une
table vibrante (shaken solutions). Cette comparaison permet de mettre en évidence l'influence de
l'histoire mécanique imposée lors du procédé de dispersion des solutions sur leur comportement
rhéologique. Le contrôle et la compréhension des phénomènes responsables des écarts de
viscosité des solutions de PEO ont été un objectif préalable à l'élaboration de membranes
chargées.
Pour des solutions de PEO de masse moléculaire moyenne de 5.106g/mol, cette disparité
de comportement rhéologique entre solutions agitées ou secouées se traduit sur la figure 39 par
une baisse de la viscosité newtonienne des solutions agitées, quelle que soit le régime de
concentration et un régime linéaire plus étendu (insert de la figure 39). Ces résultats montrent la
présence d'objets moléculaires de plus petite taille dans les solutions agitées au barreau
magnétique. Afin de déterminer la nature de ces objets, leur masse moléculaire moyenne a pu être
mesurée via la viscosité intrinsèque en utilisant l'équation d'Houwink-Mark-Sakurada (HMS)
suivante    KM  . La masse moléculaire moyenne Mw=107g/mol des objets présents dans les
46
solutions secouées correspond à la formation d'agrégats de chaînes. En revanche, les solutions
agitées contiennent des objets de masse moyenne Mw=1,7.106g/mol, plus faible que la valeur
théorique. Ces objets sont donc vraisemblablement des chaînes de polymères rompues sous
l'effet de l'élongation imposée par l'agitation, voire des agrégats de chaînes rompues. Les liaisons
rompues dans les groupes d'éthylène donnent naissance à des radicaux libres qui peuvent aboutir
à la formation de groupes - OH ; - CH3 ;-CH = CH2 ou – CH2 = CH3, selon le point de rupture
dans la chaîne. Ces trois derniers groupes ont un caractère hydrophobe qui favorise les
interactions associatives entre les chaînes rompues. Ces interactions hydrophobes expliquent la
valeur élevée des interactions de paire mesurées pour les solutions agitées et peuvent favoriser la
formation d'agrégats.
C = 1.5 wt%
2
 Pa.s
10
1
10
0
10
Shaked solution
Stirred solution
-4
10
-3
10
-2
10
-1
10
0
10
 / s1
1
10
2
10
3
10
Fig. 39: viscosité newtonienne des solutions agitées ()
et secouées () en fonction de la concentration. Insert:
comportement visqueux à la concentration de 1,5% en
masse
Fig. 40: viscosité réduite et viscosité inhérentes des
solutions agitées (, ) et secouées (, ) en
fonction de la concentration.
La dynamique moléculaire du système a été étudiée par des mesures de viscoélasticité
linéaire. L'ajustement des mesures par un modèle de Maxwell Généralisé a permis de suivre la
dynamique de relaxation des plus gros objets moléculaires, associée au temps de relaxation le plus
long.
Dans le cas des solutions secouées, ce temps de relaxation est long et il augmente avec la
concentration en polymère. (Fig. 41) Il reflète la dynamique des agrégats freinée par la présence
de chaînes non dégradées. En revanche,
pour les solutions agitées, ce temps de
relaxation augmente en régime semi
dilué puis diminue en régime concentré.
L'agitation induit par le barreau
magnétique a deux effets antagonistes:
- il favorise la formation de chaînes
courtes ayant un caractère
hydrophobe qui contribuent à leur
agrégation,
- les agrégats sont sensibles au
cisaillement et peuvent alors se
rompre
lorsque
leur
taille
augmente.
Fig. 41: Temps de relaxation des objets les plus gros présents
dans les solutions agitées () et secouées () en fonction de
la concentration
47
Le phénomène d'agrégation de chaînes courtes semble être dominant dans le régime semidilué alors que la rupture des agrégats l'emporte sous l'effet de forces hydrodynamiques
croissantes en régime concentré.
Cette influence de l'histoire mécanique des solutions sur leurs propriétés rhéologiques est
fortement dépendante de la masse
moléculaire du polymère. Ainsi, pour une
masse moléculaire plus faible (106g/mol),
seule la rhéologie des solutions en régime
concentré est sensible à l'histoire
mécanique et cette influence n'est plus
visible pour des masses moléculaires
inférieures à 3.105g/mol. Les objets
moléculaires présents en régime dilué ont
alors une masse très proche de la masse
théorique d'une chaîne de polymère. En
effet, le taux d'élongation nécessaire pour
rompre une chaîne de PEO est
proportionnel à M-2,25.19 Pour ces
Fig. 42: viscosité newtonienne des solutions agitées
polymères
de
faibles
masses
(symboles pleins) et secouées (symboles vides) en fonction
moléculaires,
les
contraintes
de la concentration pour des polymères de masse
élongationelles renforcées par des
moléculaire 106 g/mol () et 3.10 5 g/mol () en
contraintes locales des chaînes voisines
solution aqueuse.
ne sont pas suffisantes pour rompre les
chaînes en régime concentré.
Les résultats de cette étude montrant la fragilité du polymère sont d'une grande importance
pour les applications industrielles impliquant des écoulements turbulents. L'histoire mécanique
est capable de modifier les structures en solution (longueur des chaînes, taille des agrégats) et par
conséquent, les propriétés rhéologiques des solutions. Pour l'élaboration des membranes, le
procédé d'obtention par coulée/évaporation reste le procédé le moins perturbant pour la
structure des chaînes. Le procédé d'extrusion impose des cisaillements et élongations de fortes
intensités capables de rompre plus efficacement les chaînes.
2.2.2.3.
Propriété de renfort de la matrice par des charges cellulosiques
Pour l'étude des composites, notre choix pour la matrice s'est porté sur le PEO de haute
masse moléculaire (5.106g/mol), conférant a priori aux composites les propriétés mécaniques les
plus élevées. L'étude des composites a porté principalement sur deux points:
19
-
L'utilisation de deux procédés de mise en forme, par coulée/évaporation et par extrusion,
pour des composites à base de fibres courtes et rigides ou whiskers de ramie à différents
taux de charge.
-
L'utilisation de charges de morphologie et de nature différentes, avec d'une part des
whiskers de sisal, coton ou ramie et d'autre part des charges longues et flexibles ou
microfibriles de sisal.
Islam MT, Vanapalli SA, Solomon MJ, Macromolecules, 2004, 37, 1023.
Nous avons considéré l'influence du procédé et du type de charge sur la morphologie, les
propriétés rhéologiques, thermiques et mécaniques des membranes.
2.2.2.3.1.
Influence du taux de charge et du procédé de mise en forme
Les films obtenus par coulée/évaporation ont servi de matériaux de référence pour cette
étude. Le taux de charge est un paramètre important dans l'élaboration d'un composite. Pour une
température supérieure à 70°C, au-delà
de la température de fusion de la
matrice Tm = 57°C, le module élastique
des composites se stabilise sur une large
gamme de températures, et ceci dès les
plus faibles taux de charge (Fig. 43). La
valeur de stabilisation du module
élastique augmente avec le taux de
charge. Cette évolution du module avec
le taux de charge au-delà de la
température de fusion est très bien
décrite par l'association, en parallèle
d'un réseau percolant de module E'R et
le milieu dispersant constitué de la
matrice et des whiskers non percolés.
Le module élastique théorique
et alors décrit par la relation suivante:
Fig. 43: Module élastique E' normalisé en fonction de la
température pour la matrice seule () et pour un taux de charge
de 3% () , 6% () 10%(), 20% () et 30% () en masse.
, avec
lorsque la fraction volumique des fibres
est inférieure à la fraction volumique
de percolation
(
) au-delà du seuil de percolation
L'effet de renfort des composites résulte donc de la formation d'un réseau de percolation
de whiskers en interaction par liaison hydrogène. Cette interprétation est en accord avec l'étude
de la microstructure, menée par observation MEB de surfaces de cryofractures, montrant une
distribution homogène des charges dans la matrice.
Un équilibre délicat doit être trouvé entre un taux de charge élevé, procurant un effet de
renfort important et le taux de charge le plus faible, garantissant une bonne conductivité de la
membrane. Notre choix s'est donc porté sur un taux de charge de 6% en masse correspondant à
la fraction volumique
4, 86%, au-dessus du seuil de percolation.
Le procédé de mise en forme par coulée/évaporation permet d'obtenir des membranes
offrant de bonnes propriétés mécaniques jusqu'à des températures de l'ordre de 100°C.
Cependant, ce procédé nécessite la solubilisation du polymère, le séchage très long des solutions
qui empêchent son exploitation à l'échelle industrielle. Nous avons donc comparé les propriétés
d'usage des membranes obtenues par coulée/évaporation avec des membranes obtenues par
extrusion, un procédé industriel standard.
48
49
Nous avons vu précédemment comment le procédé de solubilisation de la matrice,
pourtant considéré comme faiblement perturbateur pour le polymère, pouvait cependant agir sur
la microstructure et le comportement
rhéologique de la matrice. L'utilisation d'un
procédé plus intrusif, tel que l'extrusion,
agit bien évidemment sur la matrice mais
également sur les charges et leur
organisation dans le composite. Ainsi, les
clichés MEB de cryofractures des
composites obtenus par extrusion montrent
une agrégation des charges. Par ailleurs, les
charges elles-mêmes voient leur taille
moyenne (longueur et diamètre) fortement
Fig. 44: distribution de la longueur des whiskers de ramie
diminuer après extrusion, (Fig. 44) sans
avant et après extrusion.
changement perceptible de leur rapport de
forme, de l'ordre de 25.
Les mesures viscoélastiques linéaires à T = 90°C de la Fig. 45, et les mesures de DMA de
la Fig. 46 menées sur des membranes obtenues par coulée/évaporation ou par extrusion
montrent l'effet perturbateur de l'extrusion sur la matrice et le composite. Pour la matrice, les
modules viscoélastiques des membranes extrudées sont plus faibles. Avec une masse moléculaire
moyenne de 8,7.105 g/mol pour la matrice extrudée et 1,7.106 g/mol pour la matrice
coulée/évaporée, la matrice après extrusion est plus fortement dégradée que la matrice
coulée/évaporée. De plus, les mesures de DMA montrent une chute plus importante du module
élastique pour la matrice extrudée au voisinage de la température de transition vitreuse Tg=-55°C.
L'extrusion diminue donc le taux de cristallinité de la matrice.
Films évaporés
Films extrudés
Fig. 45: Modules viscoélastiques du
composite (symboles carrés) et de sa matrice
(symboles ronds)
obtenus
a)
par
coulée/évaporation b) par extrusion.
Fig. 46: Module élastique E' normalisé pour a)
la matrice coulée/évaporée () extrudée ()
ou et b) le composite extrudé () et
coulé/évaporé ().
En ce qui concerne les composites, le comportement viscoélastique des films obtenus par
coulée/évaporation est typique d'un solide élastique avec des modules peu dépendant de la
fréquence et le module G' bien supérieur au module G", confirmant ainsi la présence du réseau
de fibres (Fig 45 a). En revanche, pour le composite extrudé, le comportement viscoélastique est
proche de celui de la matrice. Le procédé d'extrusion ne favorise pas la formation du réseau de
whiskers pour ce taux de charge. L'agrégation des charges augment artificiellement le taux de
charge nécessaire pour créer un réseau percolant.
Outre l'effet d'agrégation, l'extrusion du composite est susceptible d'orienter les fibres
courtes. Pour s'en assurer, des mesures de DMA ont été réalisées selon deux directions dans le
composite: dans la direction de l'extrusion (Fig. 46 b, ) et perpendiculairement à l'extrusion
(Fig. 46 b, ). Les modules élastiques mesurés dans ces deux directions se superposent: les
propriétés mécaniques des composites extrudés sont isotropes. L'extrusion n'oriente donc pas les
whiskers. Au-delà de la température de fusion, les modules du composite extrudé marquent un
pseudo-palier à un niveau plus faible que pour le composite obtenu par coulée/évaporation et le
module élastique décroit faiblement avec la température. En effet, la viscosité élevée de la
matrice, la cinétique rapide du procédé d'extrusion et les contraintes mécaniques élevées réduisent
la densité de liaison hydrogène entre whiskers et conduit à la formation d'un réseau faible.
Il apparaît clairement que le procédé d'extrusion dégrade la matrice, coupe et agrège les
charges de cellulose en évitant la formation d'un réseau dense de whiskers. Ce procédé, bien
qu'adapté à l'industrialisation, ne permet pas un renfort comparable à celui observé avec les
membranes obtenues par coulée/évaporation.
2.2.2.3.2.
Influence de la morphologie et de la nature des charges
Comme nous avons pu le voir, le procédé d'extrusion ne permet pas d'optimiser le renfort
mécanique à haute température. L'étude de l'influence des charges s'est donc portée sur les
membranes obtenues par coulée/évaporation. L'utilisation de whiskers de ramie, sisal et coton a
permis d'étudier l'effet de la nature des charges et la comparaison entre whiskers et microfibrilles
de sisal a permis de considérer l'effet de la morphologie des charges. Le taux de charge est de 6%
en masse pour toutes les charges utilisées. Comme le montre le tableau ci-dessous, ces charges se
distinguent par leurs rapports de forme, leurs surfaces spécifiques et leurs densités de charges de
surfaces.
Longueur L
(nm)
Diamètre d
(nm)
Rapport de
forme L/d
Surface
spécifique
(m2/g)
Densité de charges
de surface
(e/nm2)
Module
élastique
(GPa)
Wiskers de sisal
Wiskers de ramie
Wiskers de coton
Microfibrilles de sisal
Tableau 1: Dimensions et charges de surfaces des whiskers de sisal, ramie, coton et microfibrilles de sisal.
Dans le souci d'optimiser également les propriétés d'usage des membranes, leurs
caractérisations ont également été menées en présence de sel de Lithium, LiTFSI. La quantité de
sel dans la membrane est donnée par le rapport molaire oxyde d'éthylène/Lithium, noté O/Li.
50
51
En l'absence de sel, le renfort des composites est maximal avec les whiskers de sisal (Fig.
47). Le module élastique à hautes températures atteint alors une valeur de 12 MPa. Ce niveau de
renfort, bien supérieur en présence de whiskers de ramie, est confirmé par le modèle de
percolation décrit en page 49. En présence de microfibrilles de sisal, le module élastique à haute
température est significativement plus faible. Le renfort ne dépend donc pas d'interactions
spécifiques entre la charge et la matrice mais essentiellement du rapport de forme et du module
élastique des whiskers.

+


-
PEO
PEO + sisal
PEO + ramie
PEO + coton
PEO + MF
Fig. 47: Module élastique réduit E' en fonction de la température pour le PEO
et les composites.
En présence du sel de Lithium, le module élastique de la matrice diminue fortement au
passage de la température de transition vitreuse (Fig. 48). Le sel de Lithium diminue le taux de
cristallinité du PEO. En concentration élevée en sel, O/Li = 12, le module de la matrice diminue
à la température de fusion mais il reste constant à une valeur de 2 MPa jusqu'à une température
de 145°C. Cette stabilisation au module élastique est liée à la formation d'interactions entre les
ions Lithium et le PEO, agissant comme des points de réticulation pour le système.
C'est encore avec les whiskers de sisal qu'un renfort important est observé à haute
température. Ce résultat indique que la présence de sel n'affecte pas la cohésion du réseau de
fibres.

+


PEO
PEO-O/Li 12
PEO-O/Li 20
PEO-O/Li 12, WS 6%
PEO -O/Li 12, MF 6%
Fig. 48: Module élastique réduit E' en fonction de la température pour le PEO
avec et sans sel et les composites chargés à 6% en masse de cellulose et un rapport
O/Li =12.
L'évolution de la conductivité ionique de l'électrolyte polymère et des composites
polymères chargés en whiskers et microfibrilles est reportée en figure 49. Pour les composites, la
conductivité la plus élevées est atteinte en présence de whiskers de sisal, avec une valeur de
3,1.10 -4 S/cm à 60°C. A haute
température, la diminution de la
conductivité des composites par
rapport à la matrice seule est liée
à la restriction de mobilité de
l'électrolyte
à
l'interface
whiskers/électrolyte et non dans
la matrice, la Tg de la matrice
restant inchangée en présence de
cellulose. En présence de
microfibrilles, les longues fibrilles
enchevêtrées
réduisent
plus
fortement la mobilité des chaînes
de PEO. Lorsque la température
Fig. 49: Conductivité ionique de l'électrolyte PEO-LiTFSI avec
baisse, la diminution de la
O/Li = 20 () et pour les composites avec 6% en masse de whiskers
conductivité est due à la
de sisal (), ramie () et coton (+) et de microfibrilles de sisal ().
croissance de la phase cristalline.
Parmi les charges considérées, les whiskers de sisal offrent d'excellentes propriétés de
renfort mécanique sans baisser significativement la conductivité du polymère. L'utilisation d'un
procédé par coulée/évaporation réduit cependant l'intérêt d'un tel composite pour l'application
aux piles Lithium-polymère.
52
53
Chapitre 3
Perspectives
L'activité de recherche que j'ai initiée il y a un an, et que je souhaite poursuivre dans les
années à venir, regroupe les deux thématiques précédemment décrites; à savoir la rhéologie des
polymères en solution et la mise en forme de polymères. Cette thématique, centrée sur le procédé
d'electrospinning ou filage de polymères sous champs électriques intenses, a pour but de former
des nanofibres de polymères fonctionnels à partir des polymères en solution. Cette nouvelle
thématique de recherche au laboratoire a reçu en 2010 le support financier du pôle Sciences de la
Matière, INGénierie, Univers et Environnement, SMINGUE ainsi que de Grenoble Institut
National Polytechnique dans le cadre du programme Bonus Qualité Recherche, BQR.
Dans le cadre des travaux déjà engagés et en collaboration avec l'équipe Chimie et
biotechnologie des oligosaccharides du CERMAV, je co-encadre la thèse de Mlle Lancuski,
débutée en septembre 2010 et portant sur l'élaboration de matériaux nanostructurés pour
l'ingénierie tissulaire. Je travaille également avec Mr Sundaray, en post-doctorat depuis octobre
2010, en collaboration avec l'équipe ELSA du LEPMI, en continuité des travaux de thèse de Mlle
D'Aprea sur l'élaboration de membranes pour piles Lithium. L'objectif est d'utiliser un réseau de
fibres de PVDF "electrospinnées" comme renfort de membranes.
Les travaux réalisés jusqu'à présent ont permis d'optimiser le procédé d'electrospinning
par la conception et la réalisation d'un ensemble expérimental performant et d'étudier la
morphologie des fibres et quelques propriétés mécaniques de polymères mis en forme par
electrospinning.
Mes perspectives de recherche pour les années à venir visent à développer cette dernière
thématique autour de trois axes:
- Le développement du procédé d'electrospinning,
- L'étude de la structuration des polymères au cours du procédé d'electrospinning,
- Le développement d'applications pour l'ingénierie tissulaire
3.1 Développement du procédé d'electrospinning
La morphologie du réseau de fibres dépend en grande partie de la conception du
collecteur. En collaboration avec le Laboratoire des Ecoulements Géophysiques et
Industriels, LEGI, nous allons concevoir des collecteurs permettant de contrôler la
morphologie des réseaux et en particulier l'alignement plus précis des fibres. Ces travaux devront
aboutir à moyen terme au dépôt de brevets avec le support de la société Floralis, filiale de l'UJF.
Par ailleurs, les paramètres pertinents pour l'electrospinning restent encore méconnus; en
particulier les propriétés rhéologiques nécessaires pour l'electrospinning de polymères en
solution. La viscosité à taux de cisaillement nul est classiquement prise en compte pour définir la
capacité d'une solution à être "electrospinnée" alors même que le fluide n'est soumis à aucun
cisaillement mais une élongation intense. Il est donc indispensable de mesurer les propriétés du
jet de polymère en élongation. Mon objectif à court terme est donc de développer une
instrumentation permettant d'accéder à la fois au taux d'élongation et aux contraintes
élongationnelles. La taille du jet, la faible consistance du fluide (concentration légèrement
supérieure à la concentration d'enchevêtrement) et la rapidité du procédé (vitesse d'émission du
jet >10 m.s-1) rendent impossible l'utilisation de capteurs de forces. Dans le cadre de ce projet, j'ai
d'ores et déjà développé l'instrumentation pour la mesure du taux d'élongation par suivi de
particules. La détermination des contraintes élongationnelles sera développée par le suivi d'une
perturbation sous la forme d'une impulsion latérale de type acoustique ou mécanique le long du
jet. L'élargissement temporel de cette impulsion est directement relié aux contraintes
élongationelles dans le jet. Ce travail expérimental se fera en collaboration avec nos collègues
spécialistes des ultrasons qui nous ont rejoints récemment au sein du Laboratoire Rhéologie et
Procédés.
3.2 Structuration des polymères au cours du procédé d'electrospinning
Les conditions d'enchevêtrement des chaînes de polymère sont nécessaires pour obtenir
des nanofibres de polymère. Sous élongation intense, l'orientation des chaînes est fortement
suspectée mais la communauté scientifique ignore encore comment se structure le polymère
pendant ce procédé. Mon objectif est de caractériser l'organisation microstructurale des chaînes
de polymère (orientation, confinement, cristallisation, relaxation, …) induit par le procédé. Dans
un premier temps, cette caractérisation sera menée sur des nanofibres préalablement déposées sur
le collecteur. Cette étude se fera par diffusion de neutrons en collaboration avec ILL et le LLB. A
plus long terme, je souhaite mener cette caractérisation directement dans le jet de polymère,
pendant le procédé. Cette caractérisation, plus délicate, pourra se faire par diffusion de rayons X
sur la ligne ID 13 possédant une taille de faisceau de l'ordre du micron. Une collaboration avec le
Department of Mechanical Engineering Technion de l'Israel Institute of Technology, Haifa, sur
des aspects simulation numérique est en préparation.
3.3 Développement d'application pour l'ingénierie tissulaire
Je souhaite développer des applications du procédé d'electrospinning vers l'ingénierie
tissulaire. Dans ce cadre, je souhaite étudier la mise au point de matrices 3D poreuses à base de
fibres biocompatibles et biorésorbables afin de les utiliser dans le domaine de la médecine
régénératrice, comme support de croissance de cellules spécifiques.
La thèse de Mlle Lancuski, débutée en septembre 2010 et destinée à la production
d'implants de biopolymères pour la régénérescence neuronale, s'inscrit très clairement
dans cet axe de recherche. Accroître la vitesse de croissance des axones de façon unidirectionnelle
(donc stimuler leur régénération et favoriser leur guidage) représente un espoir réaliste de
guérison de certaines neuropathologies. Une voie prometteuse pour atteindre cet objectif consiste
à mettre en œuvre par electrospinning une structure "support" poreuse de nanofibres de
polymères fortement alignées, servant de guide pour une croissance optimisée des axons. La
maîtrise de la morphologie du réseau de fibres et le suivi de la croissance des neurones dans le
réseau sont des points déterminants dans le succès de ce projet. Des observations de la croissance
des neurones sur le réseau de fibres par tomographie X sur la ligne ID 22 et par spectrométrie de
corrélation de fluorescence en association avec de la microscopie confocale seront programmées.
Une stratégie pour faciliter la croissance des neurones consiste à préparer une matrice
biocompatible incorporant des ligands capables de promouvoir l’adhésion et la croissance des
neurones. Les oligosaccharides tels que le HNK1 font partie de ces ligands potentiels. Ce projet à
long terme est mené en collaborations avec Sébastien Fort et Bernard Priem de l'équipe
Glycochimie du CERMAV, UPR 5301, pour l'aspect fonctionnalisation du réseau et Fatia
Nothias et Sylvia Soares de l'équipe Régénération et Croissance de l'Axone du Laboratoire
Physiopathologie des Maladies du Système nerveux Central, INSERM UMRS-952, CNRS UMR
7224, Université Pierre et Marie Curie, Paris VI, pour le volet croissance cellulaire.
54
55
Un second domaine d'application serait l'élaboration d'artères de petits diamètres
(< 6 mm) pour permettre la croissance de cellules endothéliales de la paroi interne du
vaisseau sanguin. Les substituts synthétiques, utilisant notamment du textile, sont performants
pour le remplacement d’artères de moyens et gros diamètres, mais sont peu satisfaisants lorsqu’il
s’agit d’artères de petits diamètres. Le procédé d'élaboration par electrospinning peut être une
alternative aux techniques existantes telle que le tissage. Un tel projet pourra être développé dans
le cadre du Carnot PolyNat dans lequel le laboratoire Rhéologie et Procédés émarge.
Enfin la troisième application visée est la réparation tissulaire de peau pour grands
brûlés. Plusieurs produits issus de travaux de recherche ont déjà été expérimentés mais sans
réelle commercialisation en France. Nous envisageons l'utilisation de polymères de type collagène
ou chitosan pour produire des "peaux artificielles".
Ces applications nécessitent de nouer des collaborations vers la recherche clinique. Dans ce
domaine, l'environnement grenoblois est particulièrement riche avec, entre autres, les équipes
Dynamiques Cellulaire / Tissulaire et Microscopie Fonctionnelle (DyCTiM) et Gestes MédicoChirurgicaux Assistés par Ordinateur (GMCAO) du laboratoire TIMC, l'Institut des
neurosciences de Grenoble et le centre de recherche biomédical CLINATEC® expert en
neurochirurgie. A l'échelle régionale, j'envisage également de collaborer avec le Laboratoire des
Substituts Cutanés - Banque de cornées -Banque de tissus et de cellules du Groupement
Hospitalier Edouard Herriot de Lyon. Le transfert technologique pourrait se faire avec l'aide de
l'Institut Français de l'habillement et du textile (IFTH), et des PME comme ABISS spécialisée
dans la conception et la fabrication d'implants chirurgicaux pour les tissus mous et Texinov,
spécialiste du textile technologique.
La croissance cellulaire, point commun à toutes ces applications, est optimisée lorsque la
structure support est en tension. La caractérisation mécanique de ces supports s'avère donc
nécessaire pour mener à bien ces projets. Ces tests de résistance mécanique seront menées en
collaboration avec l'équipe Mécanique et Couplages Multiphysiques des Milieux Hétérogènes,
(CoMHet), du laboratoire Sols, Solides, Structures - Risques (3SR). Un couplage entre les
propriétés mécaniques et la microstructure mentionnée précédemment sera particulièrement
pertinent.
Synthèse
Expériences professionnelles
Maître de conférences (laboratoire de rhéologie et IUT)
2 post-doctorats
2 postes d'ATER
Publications
17 articles internationaux de rang A
1 chapitre d'ouvrage
17 conférences internationales
3 congrès nationaux
Encadrements
1 Post-doctorat
2 Thèses
1 PFE
7 Masters
Responsabilités collectives
10 revues d'articles pour des journaux internationaux
Membre élu au conseil du GFR
Représentant élu du personnel au laboratoire de rhéologie, Grenoble
Membre élu au conseil du département GMP de l'IUT
Membre de 2 comités de sélection de maîtres de conférences
Co-organisateur du congrès international "Annual Alpine Rheology
Meeting"
Responsable du site du GFR
Responsable des poursuites d'études à l'IUT, département GMP
Enseignements
IUT 1 Grenoble, département Génie Mécanique et Productique
Mécanique, mathématiques
UFR de chimie, Génie des systèmes industriels
Rhéologie
Volume horaire: ~260h eq. TD annuel
Thèmes de Recherche
Rhéologie des polymères associatifs
Mise en forme de polymères
Projet: Mise en forme de nanofibres par electrospinning
Développement du procédé d'electrospinning
Structuration des polymères au cours du procédé
d'electrospinning
Développement d'applications pour l'ingénierie tissulaire
56
57
Annexes
Aubry, T., F. Bossard, G. Staikos, G. Bokias, "Rheological study of semi-dilute aqueous
solutions of a thermoassociative copolymer", J. Rheol., 47(2); 577-587, 2003.
Bossard, F., V. Sfika, C. Tsitsilianis "Rheological Properties of Physical Gel formed by Triblock
Polyampholyte in Salt-Free Aqueous Solutions", Macromolecules, 37, 3899-3904, 2004.
Bossard, F., M. Sotiropoulou and G. Staikos, "Thickening effect in Soluble Hydrogen-bonding
Interpolymer complexes. Influence of molecular composition and pH" J. Rheol., 48(4),
927-926, 2004.
Bossard, F., V. Sfika, C. Tsitsilianis and S. Yannopoulos, "A Novel Thermothickening
Phenomenon Exhibited by a Triblock Polyampholyte in Aqueous Salt-Free Solutions",
Macromolecules, 38, 2883-2888, 2005.
Bossard, F., T. Aubry, G. Gotzamanis, C. Tsitsilianis, "pH-tunable Rheological Properties of a
Telechelic Cationic Polyelectrolyte Reversible Hydrogel" Soft Matter, 2, 510-516, 2006.
Sotiropoulou, M, F. Bossard, E. Balnois, J. Oberdisse and G. Staikos, "Characterization of the
Core-Shell Nanoparticles Formed as Soluble at low pH Hydrogen bonding Interpolymer
Complexes", Langmuir, 23, 11252 –11258, 2007.
Bossard, F., M. Moan, T. Aubry, "Linear and non-linear Viscoelastic Behavior of Very
Concentrated Kaolinite Suspensions", J. Rheol. 51, 1253-1270, 2007.
Bossard, F., I. Pillin, T. Aubry and Y. Grohens, "Rheological Characterization of Starch
Derivatives/Polycaprolactone Blends Processed by Reactive Extrusion", Polym. Eng. Sci,
48, 1862-1870, 2008.
Bossard, F., N. El Kissi, A. D'Aprea, F. Alloin, J-Y Sanchez and A. Dufresne, " Influence of
dispersion procedure on rheological properties of aqueous solutions of high molecular
weight PEO", Rheologica Acta. 49, 529–540, 2010.
Alloin F. , A. D’Aprea, A. Dufresne, N. El Kissi, F. Bossard, " Nanocomposite polymer
electrolyte based on whisker or microfibrils polyoxyethylene nanocomposites",
Electrochimica Acta, 55, 5186–5194, 2010.
Alloin F. , A. D’Aprea, A. Dufresne, N. El Kissi, F. Bossard, "Poly(oxyethylene) and ramie
whiskers based nanocomposites. Influence of processing: extrusion and
casting/evaporation", Cellulose, 2011.
Rheological study of semidilute aqueous solutions
of a thermoassociative copolymer
Thierry Aubrya) and Frédéric Bossard
Laboratoire de Rhéologie, Université de Bretagne Occidentale, 6 avenue Victor Le
Gorgeu, 29285 Brest Cedex, France
G. Staikos and G. Bokias
Department of Chemical Engineering, University of Patras and Institute of
Chemical Engineering and High Temperature Processes, ICE/HT-FORTH,
P.O. Box 1414, 26504 Patras, Greece
(Received 15 October 2002; final revision received 11 December 2002)
Synopsis
In this paper, the linear and nonlinear rheological behavior of semidilute aqueous solutions of a
recently synthesized thermoassociative graft copolymer was investigated, as a function of
temperature and polymer concentration. The polymer, namely CMC–g–PNIPAM, is based on a
carboxymethylcellulose 共CMC兲 backbone bearing thermosensitive poly共N-isopropylacrylamide兲
共PNIPAM兲 sidechains. The samples have been submitted to steady shear, oscillatory shear, and
step-strain experiments, mainly at temperatures above the threshold temperature T assoc to observe
thermothickening. The linear and nonlinear rheological data clearly show the existence of two
temperature regimes, separated by a transition temperature T ⬘ ⬎ T assoc . At temperatures below
T ⬘ , the solutions behave like a soft critical gel, corresponding to weak PNIPAM segregation. At
temperatures above T ⬘ , the solutions behave like a stiff critical gel, corresponding to strong
PNIPAM segregation. © 2003 The Society of Rheology. 关DOI: 10.1122/1.1545077兴
I. INTRODUCTION
Nowadays water-soluble associating polymers are extensively used as rheology modifiers, mainly as thickeners, in numerous industrial applications 关Glass 共1989兲; Shalaby
et al. 共1991兲; Aubry and Moan 共1997兲兴. These complex fluids, like most polymer solutions and simple fluids, exhibit a thermothinning behavior, characterized by a decrease of
the viscosity when temperature is increased, usually described by the Arrhenius law.
Thermothickening polymeric systems constitute a relatively new class of water-soluble
associative polymers which exhibit an increase of the viscous properties as the temperature is increased. This peculiar thermal behavior is due to the presence of reversible
intermolecular associations that are favored on warming 关Sarrazin-Cartalas et al. 共1994兲;
Loyen et al. 共1995兲; Bokias et al. 共1997兲兴; they are of great potential interest in the
formulation of fluids submitted to heat. Among these materials, water-soluble thermoresponsive graft copolymers, obtained by grafting polymer chains, with a lower critical
solution temperature 共LCST兲 in water, onto a hydrophilic backbone, have been shown to
a兲
Author to whom all correspondence should be addressed; electronic mail: [email protected]
© 2003 by The Society of Rheology, Inc.
J. Rheol. 47共2兲, 577-587 March/April 共2003兲
0148-6055/2003/47共2兲/577/11/$25.00
577
578
AUBRY ET AL.
form a particularly promising class 关Hourdet et al. 共1997, 1998兲; Durand and Hourdet
共1999, 2000兲兴. The design of these thermoassociative polymeric systems is based on the
limitation to a microscopic level of the phase separation of the LCST grafts above a
critical temperature T assoc . T assoc is the threshold temperature to form a reversible associative network of physically crosslinked hydrophilic backbones connected via hydrophobic graft aggregates 关Hourdet et al. 共1998兲兴. The value of T assoc is usually equal to the
LCST of the graft 关Hourdet et al. 共1998兲兴, still it may be slightly different in the case
where the hydrophilic backbones are stiff chains, which induce topological constraints on
grafts and perturb their phase separation 关Bokias et al. 共2001兲兴.
Contrary to thermothinning associative polymers which have been widely and thoroughly studied from a rheological point of view over the last decade 关Annable et al.
共1993兲; Aubry and Moan 共1994兲; Plazek and Frund 共2000兲; Berret et al. 共2001兲兴, thermoassociative polymers have not received much attention from rheologists. Indeed most
rheological data reported in the literature only concern the linear and nonlinear viscous
properties of thermothickening polymeric solutions. In the present work, we report on a
more complete rheological characterization of a thermoassociative water-soluble polymer, drawing more specific attention to the linear and nonlinear viscoelastic properties of
the reversible thermoassociative network formed above the threshold temperature T assoc .
The rheological investigation has been performed with aqueous solutions of a recently
synthesized thermoassociative graft copolymer, CMC–g–PNIPAM, based on a carboxymethylcellulose
共CMC兲
backbone
bearing
thermosensitive
poly共Nisopropylacrylamide兲共PNIPAM兲 sidechains. This thermothickening polymeric material is
attractive because of the high biodegradability of the CMC backbone and also because its
thermothickening properties have been shown to be effective over a very large pH region
关Bokias et al. 共2001兲兴.
II. EXPERIMENT
Sodium CMC used in this work has a weight average molecular weight of 8.2
⫻104 g/mol and an intrinsic viscosity of about 165 ml/g. Amino-terminated PNIPAM
chains, prepared by radical polymerization, with an average molecular weight of 4.3
⫻104 g/mol and an intrinsic viscosity of 32 ml/g, were grafted onto the CMC backbone
in water, to get the CMC–g–PNIPAM copolymer 关Bokias et al. 共2001兲兴. The thermoassociative character of this polymer is due to the high thermal sensitivity of the PNIPAM
chains. Indeed turbidity measurements clearly proved that PNIPAM phase separates from
aqueous solution at a temperature of 33 °C 关Bokias et al. 共2001兲兴.
The sample synthesized for the study is estimated to contain one PNIPAM chain per
one copolymer chain, in average. However, the actual distribution of the sidechains on
the CMC backbones remains unknown and there is certainly some amount of CMC
chains bearing more than one PNIPAM chain.
The results reported in this paper concern CMC–g–PNIPAM solutions in pure water
(pH ⬃ 8). Under these conditions, the polymer chains are considerably expanded due to
electrostatic repulsions between the carboxylate groups of the CMC chain. All polymer
solutions were prepared by dissolving the appropriate amount of polymer in pure water
under gently stirring for 24 h. All the solutions prepared remained clear on the whole
range of temperatures imposed, showing that no macroscopic phase separation ever occured.
Four concentrations have been studied: 3%, 4%, 5%, and 6% w/w polymer solutions.
Actually the relevant concentration to be considered in this work is the relative concentration of the CMC backbones, to be compared with the overlap concentration c * and the
SOLUTIONS OF A THERMOASSOCIATIVE COPOLYMER
579
FIG. 1. Apparent viscosity as a function of shear stress, for a 6% w/w polymer solution at various temperatures.
Open symbols correspond to creep measurements and full symbols to steady shear measurements.
entanglement concentration c e of CMC backbones in water. In our study, the relative
concentration of the CMC backbones lies between 2% and 4% w/w, which is higher than
c * ⬃ 0.5% w/w and higher than c e ⬃ 1% w/w, so that CMC backbones were assumed
to be in a semidilute entangled regime. Besides, the threshold temperature T assoc of these
semidilute CMC–g–PNIPAM solutions has been shown to be 35 °C, that is slightly
higher than the LCST of the PNIPAM chains in pure water 关Bokias et al. 共2001兲兴.
The linear and nonlinear rheological properties of the polymer solutions were studied
using two rheometers: 共i兲 a Rheometric Scientific ARES, equipped with either a cone and
plate geometry (radius ⫽ 25 mm, cone angle ⫽ 2.3°) or a Couette geometry (gap
⫽ 1 mm) and 共ii兲 a Carri-Med CSL50 constant stress rheometer, equipped with a cone
and plate geometry (radius ⫽ 30 mm, cone angle ⫽ 2°). The imposed temperature
ranges from 30 °C up to 55 °C. To prevent dehydration from the solution, a thin layer of
low-viscosity silicone oil was put on the air/sample interface.
III. RESULTS
A. Flow curves
Figure 1 shows the flow curve of a 6% w/w CMC–g–PNIPAM solution in water at
various temperatures. These results are quite representative of those obtained for any
solutions tested. All samples exhibit a low-shear Newtonian plateau, determined by creep
tests performed on the constant stress rheometer, whose viscosity level strongly increases
with temperature. The thermothickening effect is very pronounced: for a 6% polymer
solution it is characterized by a nearly 6 decades increase of viscosity when temperature
is increased from 30 to 55 °C.
Figure 2 shows the variation of the zero-shear viscosity as a function of temperature
for the four CMC–g–PNIPAM solutions studied. This result shows that thermothickening properties are significantly enhanced by increasing polymer concentration. Moreover,
the thermothickening regime can be divided into two temperature regimes.
580
AUBRY ET AL.
FIG. 2. Zero-shear viscosity as a function of temperature, for 3%, 4%, 5%, and 6% w/w CMC–g–PNIPAM
solutions.
共1兲 A ‘‘low’’ temperature regime, where the Newtonian viscosity ␩ 0 increases exponentially with temperature T according to the law: ␩ 0 ⫽ K(c)e T , where K(c) is a
parameter depending on the polymer concentration c.
共2兲 A high temperature regime, where the Newtonian viscosity ␩ 0 increases exponentially with temperature T according to the law: ␩ 0 ⫽ K ⬘ (c)e ␣ (c)T , where K ⬘ (c)
and ␣ (c) ⬍ 1 are parameters depending on the polymer concentration c. ␣ decreases with increasing polymer concentration: ␣ ⫽ 0.65 at c ⫽ 3%; ␣ ⫽ 0.25 at
c ⫽ 4%; ␣ ⫽ 0.15 at c ⫽ 5%; ␣ ⫽ 0.1 at c ⫽ 6%.
The transition temperature T ⬘ , marking the passage from one regime to the other, is
higher than T assoc and slightly decreases as the polymer concentration increases: T ⬘
⬃ 45 °C at c ⫽ 3%; T ⬘ ⬃ 43 °C at c ⫽ 4%; T ⬘ ⬃ 41 °C at c ⫽ 5%; T ⬘ ⬃ 40 °C at
c ⫽ 6%.
As far as the non-Newtonian behavior is concerned, it has to be noticed first that, as
the solution viscosity increases upon heating, the onset of nonlinear behavior appears at
lower and lower stresses, as shown in Fig. 1. Significant nonlinear response, namely
shear-thinning behavior, appears at temperatures above T assoc . In the ‘‘high’’ temperature
regime, that is above T ⬘ , the Newtonian plateau is followed by a weak shear-thickening
behavior, followed by a drastic shear-thinning region. The dramatic decrease of viscosity
above a certain shear stress may even be seen as the signature of an apparent yield stress
关Barnes 共1998兲兴.
B. Oscillatory shear
The storage modulus G ⬘ and loss modulus G ⬙ are plotted in Figs. 3 and 4, respectively, as a function of the shear strain at a frequency of 1 Hz, for a 6% w/w CMC–g–
PNIPAM solution. These results are quite representative of those obtained for any solutions tested.
From the results shown it occurs that both viscoelastic moduli are significantly enhanced on heating, and that the sample response is more elastic than viscous, at least at
581
FIG. 3. Storage modulus G ⬘ as a function of shear strain, at a frequency of 1 Hz, for a 6% w/w polymer
solution.
the frequency of 1 Hz. Besides, the extent of the linear viscoelastic regime decreases as
the temperature increases.
Let us now look first at the linear viscoelastic response to oscillatory shear. Figures 5
and 6 show the plateau storage modulus and the plateau loss modulus, respectively, as a
function of temperature, at the frequency of 1 Hz, for the four polymer concentrations
tested. Let us stress that the term ‘‘plateau’’ is used in this paper to qualify the strain
independent 共i.e., linear兲 viscoelastic moduli determined at a fixed frequency. These plots
show that the linear viscoelastic characteristics are enhanced by polymer addition. Moreover, the two temperature regimes defined earlier from the zero-shear viscosity measurements, are also observed in the linear viscoelastic data: the zero-shear Newtonian viscosity and the linear viscoelastic moduli are material properties which exhibit the same
FIG. 4. Loss modulus G ⬙ , at a frequency of 1 Hz, as a function of shear strain, for a 6% w/w polymer solution.
582
AUBRY ET AL.
FIG. 5. Plateau storage modulus G 0⬘ , at a frequency of 1 Hz, as a function of temperature, for 3%, 4%, 5%, and
6% w/w CMC–g–PNIPAM solutions.
qualitative dependence as a function of temperature. The same quantitative temperature
dependence is even observed below T ⬘ for these three material properties: ␩ 0 ,G ⬘0 ,G ⬙0
⬇ e T . Above T ⬘ the temperature dependence is described by the function e ␣ (c)T , ␣ (c)
decreases when the polymer concentration is increased: 共i兲 for G ⬘0 : ␣ ⫽ 0.13 at c
⫽ 3%; ␣ ⫽ 0.1 at c ⫽ 4%; ␣ ⫽ 0.09 at c ⫽ 5%; ␣ ⫽ 0.07 at c ⫽ 6% and 共ii兲 for
G 0⬙ : ␣ ⫽ 0.08 at c ⫽ 3%; ␣ ⫽ 0.06 at c ⫽ 4%; ␣ ⫽ 0.04 at c ⫽ 5%; ␣ ⫽ 0.03 at
c ⫽ 6%.
As far as the nonlinear viscoelastic response to oscillatory shear is concerned, Fig. 4
shows that, above the transition temperature T ⬘ , the sample exhibits an extra viscous
FIG. 6. Plateau loss modulus G ⬙0 , at a frequency of 1 Hz, as a function of temperature, for 3%, 4%, 5%, and
6% w/w CMC–g–PNIPAM solutions.
583
FIG. 7. Stress relaxation function vs time, for a 4% w/w CMC–g–PNIPAM solution at 37 °C.
dissipation, characterized by a hump in the G ⬙ curve. The G ⬙ hump lies in the weakly
nonlinear regime and appears at temperatures above T ⬘ , and it is all the more marked as
the temperature is higher. From a phenomenological point of view, this behavior has
some similarity with the shear-thickening behavior observed in the weakly nonlinear part
of the flow curve, at temperatures above T ⬘ .
C. Stress relaxation
All samples have been submitted to step-strain experiments, yielding the stress relaxation function, in order to study the time-dependent viscoelastic behavior. The results are
plotted in Fig. 7 for a 4% w/w CMC–g–PNIPAM solution at 37 °C, and in Fig. 8 for a
FIG. 8. Stress relaxation function vs time, for a 6% w/w CMC–g–PNIPAM solution at 45 °C.
584
AUBRY ET AL.
6% w/w CMC–g–PNIPAM solution, at 45 °C. Figures 7 and 8 are quite illustrative of
what is obtained with all solutions tested at temperatures below T ⬘ and above T ⬘ , respectively.
1. At temperatures below T ⬘
In the linear and nonlinear viscoelastic regime, the decay of the stress relaxation
function G(t) can be adequately fitted by a power-law function of time G(t) ⬇ t ⫺⌬ ,
with a characteristic exponent ⌬ of about 0.7, on the whole range of polymer concentration explored.
2. At temperatures above T ⬘
The linear viscoelastic response has a different temperature dependence, compared to
the nonlinear one. In the linear viscoelastic regime, the decay of the stress relaxation
function G(t) can be adequately fitted by a power-law function of time, with a characteristic exponent ⌬ of about 0.2, on the whole range of polymer concentration explored.
At higher strains, in the weakly nonlinear viscoelastic regime, precisely where the G ⬙
hump is observed, the nonlinear stress relaxation function exhibits a two-step response: at
short times the relaxation modulus is constant, then the system relaxes according to a
power-law decay, with the same exponent as that of the linear stress relaxation function.
At still higher shear strains, the decay of the relaxation function can be fitted by a
power-law, with nearly the same exponent as that of the linear relaxation function.
IV. DISCUSSION
A. Linear behavior
At temperatures above the association temperature T assoc , PNIPAM sidechains selfaggregate, forming hydrophobic microdomains which connect CMC chains into a physical three dimensional transient network 关Hourdet et al. 共1998兲; Durand and Hourdet
共1999兲兴. A rheological signature of the network formation is the shear-thinning behavior
observed at T ⬎ T assoc , whereas the behavior is essentially Newtonian at T
⬍ T assoc .
From a phenomenological point of view, and following the classical theory of transient
networks 关Green and Tobolsky 共1946兲兴, the increase of the number and/or ‘‘strength’’
共characteristic time scale兲 of the intermolecular physical crosslinks could explain the
increase of the zero-shear viscosity ␩ 0 共idem for the plateau modulus G 0⬘ and loss
modulus G ⬙0 ) as the temperature and/or polymer concentration increases. From a molecular point of view, an increase of the number and/or strength of crosslinks is due to an
increase of the number of PNIPAM participating to hydrophobic clusters and/or an increase of the PNIPAM concentration within the micelles 关Hourdet et al. 共1998兲兴. At last
it should be stressed that the role played by the copolymer chains bearing more than one
PNIPAM sidechain in the formation of the transient network is paramount: they are the
elastically active chains of the network.
The enhancement of the thermothickening material properties of the CMC–g–
PNIPAM solutions, as temperature and/or polymer concentration is increased, depends on
the temperature range: it is marked in the ‘‘low’’ temperature regime, and less pronounced in the ‘‘high’’ temperature regime, where it even seems to level off as polymer
concentration increases. We think that the existence of the two temperature regimes,
separated by T ⬘ ⬎ T assoc , expresses the existence of two segregation regimes, as proposed by Hourdet and co-workers from thermodynamic considerations 关Hourdet et al.
共1998兲兴.
585
• The low temperature regime would correspond to a weak segregation regime, characterized by weak hydrophobic interactions, leading to the formation of loose PNIPAM
aggregates which can be reinforced continuously as temperature and/or polymer concentration increases.
• The high temperature regime would correspond to a strong segregation regime, where
hydrophobic interactions and PNIPAM aggregates are strong. In this regime, increasing
temperature and/or polymer concentration is less effective than in the previous regime:
increasing the number and/or strength of strong crosslinks in a strong network is less
effective than increasing the number and/or strength of loose crosslinks in a weak network.
This schematic picture is confirmed by the linear shear relaxation data, showing that
the system relaxes according to a power-law function of time G(t) ⬇ t ⫺⌬ , whose decay
is more rapid in the low temperature regime than in the high temperature one. Indeed
such relaxation patterns indicate that the thermoassociative polymer solutions behave like
soft critical gels, characterized by a high network specific exponent ⌬, below T ⬘ ;
whereas they behave like stiff critical gels, characterized by a low network specific
exponent ⌬, above this transition temperature 关Winter and Mours 共1997兲兴.
B. Nonlinear behavior
The qualitative features of the nonlinear part of the flow curves resemble those obtained with hydrophobically associating polymeric systems 关Aubry and Moan 共1994兲;
Volpert et al. 共1996兲兴. This rheological similarity is quite expected as both thermoassociative and classical hydrophobically associative polymer solutions are composed of a
transient physical network, through weak reversible associations, at least in the semidilute regime.
In the weakly nonlinear regime, the shear-thickening behavior in the flow curve and
the strain hardening behavior, associated with the G ⬙ hump in the viscoelastic response
has been observed with other associating systems 关Annable et al. 共1993兲; Vermant et al.
共2000兲兴. Though the physical origin of these nonlinear behaviors is still a matter of
debate, we think that they may be due to the stretching of the PNIPAM chains, following
recently published interpretations of the rheology of associative polymeric systems 关Tirtaatmadja et al. 共1997兲兴. The stretch of the PNIPAM chains could also explain the constant short time response of the stress relaxation function in the weakly nonlinear regime.
Strong enough junctions are needed to allow PNIPAM chains to stretch without breaking
the network, so that these weakly nonlinear phenomena appear in the high temperature
regime, where segregation is strong.
The strong nonlinear behavior, as exhibited by the drastic shear-thinning in the flow
curve and the rapid decrease of the viscoelastic moduli at high strains, is certainly the
result of the destruction of the associative network at shear rates higher than the disengagement rate from a hydrophobic junction 关Aubry and Moan 共1994兲兴.
V. CONCLUDING REMARKS
The aim of the present paper was to study the temperature and polymer concentration
effects on the linear and nonlinear rheological properties of semidilute aqueous solutions
of a thermoassociative graft copolymer.
From a phenomenological point of view, the whole set of rheological data led us to
determine a transition temperature T ⬘ , higher than the threshold temperature T assoc to
observe thermothickening, which defines the existence of two temperature regimes. At
586
AUBRY ET AL.
temperatures below T ⬘ , the solutions behave like a soft critical gel, whereas they behave
like a stiff critical gel at temperatures above T ⬘ .
From a more microscopic point of view, these results have been interpreted in terms of
connectivity of the transient associative network
共1兲 Below T ⬘ , the network junctions are weak because PNIPAM chains phase separate in
a weak segregation regime; and
共2兲 above T ⬘ , the network junctions are strong because PNIPAM chains phase separate
in a strong segregation regime.
Of course, we are quite aware that this rheological study gives only slight physical
insight into the microstructure of this complex polymeric system. Indeed we would like
to stress that our approach is essentially phenomenological in nature, so that interpretation at the molecular level is actually somewhat speculative. Additional physical insight
would be gained by the study of the influence of changes of molecular parameters 共e.g.,
graft length and number, molecular weight...兲 on the rheological properties of the polymer solutions. Moreover other physical investigation techniques, e.g., neutron scattering
techniques, would certainly be useful in order to confirm or qualify our molecular interpretations, which have been derived from a macroscopic approach. Indeed some important questions remain open:
共1兲 What is the average number and concentration of PNIPAM chains in a micellar
junction, at different temperatures and concentrations?
共2兲 Which are the effects of entanglements in the dynamics of these solutions?
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Macromolecules 2004, 37, 3899-3904
3899
Rheological Properties of Physical Gel Formed by Triblock
Polyampholyte in Salt-Free Aqueous Solutions
Frédéric Bossard,‡ Vasiliki Sfika,† and Constantinos Tsitsilianis*,†,‡
Department of Chemical Engineering, University of Patras, 26500 Patras, Greece, and Institute of
Chemical Engineering and High-Temperature Chemical Processes, ICE/HT-FORTH, P.O. Box 1414,
26504 Patras, Greece
Received September 17, 2003; Revised Manuscript Received March 18, 2004
ABSTRACT: Linear and nonlinear viscoelastic properties of an asymmetric triblock copolymer poly(acrylic
acid)-poly(2-vinylpyridine)-poly(acrylic acid) (PAA-P2VP-PAA) in salt-free aqueous solutions have been
investigated. At pH 3.4, long-range electrostatic interactions prevail, due to protonated P2VP units and
negative PAA end groups. Above a critical Cg ) 2.5% w/w, a transient network is formed through
intermolecular electrostatic interactions between negatively charged groups located at the end PAA blocks
and positively charged protonated pyridine groups located at the middle long P2VP block. The so-formed
network exhibits some atypical rheological behavior characterized by a strain hardening of storage modulus
in intermediate strain amplitudes and a pronounced shear thickening in moderated shear stresses. The
shear-induced changes in the structure of the network have been attributed to enhancement of the number
of elastically active bridges through association of free dangling ends and a transition from intra- to
intermolecular association.
Introduction
Associative water-soluble polymers are polymers that
self-assemble through temporary junctions of functional
groups. Through this generic description, different
varieties of associative thickeners can be distinguished.
Traditional associative polymers consist of neutral
polymers containing hydrophobic associative groups. In
this category, telechelic polymers (chains end-capped by
hydrophobic groups) are often considered as a model of
hydrophobic associative polymers. Extensive investigations of telechelic solutions (mostly PEO polymer derivatives) using light scattering and rheometric techniques have been carried on during this past decade.
Above a critical association concentration, polymer
chains self-associate in flowerlike micelles constituted
by a hydrophobic core, surrounded by a corona of
hydrophilic polymer loops.1 With increasing concentration, a second association process occurs between flower
micelles, corresponding to the formation of bridges.2-4
The interconnections of flowerlike micelles leads generally to a sol-gel transition which produces a sharp
increase of the viscosity. The nonlinear behavior of
telechelic associative polymers exhibits a local shear
thickening proceeding a pronounce shear thinning. In
the linear viscoelastic range a Maxwell model, corresponding to a single relaxation time and an elastic
plateau modulus, generally describes dynamic moduli.
Charged polymers display also an associative character and can be classified in two typical groups according to their architecture: hydrophobically associative
polyelectrolytes and polyampholytes. In the former
group, telechelic polyelectrolyte like its neutral parents
is end-capped by hydrophobic groups, but the ionic
character of the main backbone induces additional
interactions such as Coulomb repulsive forces between
†
University of Patras.
ICE/HT-FORTH.
* Corresponding author: Fax +30 2610 997 266; e-mail
[email protected].
‡
monomers (intramolecular) and also between chains
(intermolecular) and interactions with counterions. In
dilute solution, polymer chains can be aggregated in
finite size clusters, resulting from equilibrium between
the energy of attraction of hydrophobic stickers and the
contribution of the additional electrostatic forces.5,6 A
direct consequence of the intramolecular repulsions is
an expected stretched conformation of the polyelectrolyte into the cluster. Above a percolation concentration,
clusters are connected by stretched polymer chains in
a transient three-dimensional network. Such particular
microstructure is responsible for the unusual rheological
behavior characterized by an apparent yield stress, low
gelation concentration, short linear viscoelastic range
close to 1%, and high plateau modulus.7,8
Instead of being localized at the ends, hydrophobic
groups can be distributed along the backbone. Hydrophobically modified alkali-swellable emulsion polymers
(HASE) can be classified in this group. While the
polymers are in the emulsion state at low pH, they
become water-soluble above pH 7 by ionization of
polymer chains. Above a critical concentration, a transient network is formed through intermolecular association of the hydrophobic groups. The rheological behavior
of HASE physical gel is somehow comparable to that of
nonionic telechelic polymers, showing a more or less
pronounced shear thickening at moderate shear rate
and a sharp increase of the Newtonian viscosity and the
plateau storage modulus with increasing concentration.
However, HASE systems present an unusual viscoelastic behavior characterized by a local increase of both G′
and G′′ moduli at intermediate strain amplitudes.10-13
Such a strain dependence of G′ and G′′ moduli seems
to be typical of HASE systems.
Polyampholytes are another category of charged
water-soluble associative polymers in which opposite
charges coexist along the macromolecular chain. In this
case electrostatic interactions between oppositely charged
units are responsible for their behavior in aqueous
media.14,15 Recently, a rich phase behavior was observed
10.1021/ma0353890 CCC: $27.50 © 2004 American Chemical Society
Published on Web 04/22/2004
3900 Bossard et al.
by a weak polyampholyte of a block copolymer architecture, i.e., poly(acrylic acid)-b-poly(vinylpyridine)-bpoly(acrylic acid) (PAA134-P2VP628-PAA134).16 Depending on pH of the solution three distinct regions can be
distinguished. In the high-pH region compact micelles
with P2VP hydrophobic cores and PAA charged chains
in the corona were formed. In the intermediate-pH
region and around the isoelectric point the polymer
precipitates. Finally, in the low-pH region and low
concentrations the polymer is molecularly dissolved. The
net charge of the polyampholyte is positive since the
majority of the P2VP units are protonated, exhibiting
a polyelectrolyte character.
A novel behavior was observed at pH 3.4 (close to the
phase separation limit) and at elevated concentrations.
Despite the lack of hydrophobic blocks (which play the
role of stickers in physical gelation phenomena), this
polymer is self-assembled to a three-dimensional transient network through electrostatic interactions. In the
present article the rheological properties of this novel
physical gel will be demonstrated and compared with
those of other hydrophobically associated water-soluble
polymers. It should be mentioned here that although
much work has been reported concerning polyampholyte
chemical gels,17 this is the first work dealing with a
polyampholyte physical gel.
Macromolecules, Vol. 37, No. 10, 2004
Figure 1. Apparent viscosity as a function of shear stress
for 2.5, 3, 3.5, 4, 4.5, and 6 wt % polymer solution in water.
Full symbols correspond to data obtained by increasing stress
and open symbols to that obtained by decreasing stress.
Experimental Section
Material and Solution Preparation. The polymer studied
is a poly(2-vinylpyridine), P2VP, end-capped by two poly(acrylic acid), PAA, chains. The desired PAA-P2VP-PAA has
been obtained by acid-catalyzed hydrolysis of poly(tert-butyl
acrylate)-poly(2-vinylpyridine)-poly(tert-butyl acrylate), PtBAP2VP-PtBA, in dioxane.16 The Mw of the copolymer is 8.5 ×
104 g/mol, corresponding to a degree of polymerization of 628
for the central P2VP block and 135 for each PAA end-blocks.
These ionizable blocks are respectively weak acid and weak
basic moieties. The copolymer is water-soluble, and its aqueous
solutions give a pH close to 3.5 All polymer solutions were
obtained by dissolving the appropriate amount of polymer in
pure water. Shaking ensured the homogenization of the
solutions. For the more concentrated solutions, this shake was
assured by a series of brief centrifugations of the samples.
Between each centrifugation, the position of the samples
rotates. The measurements were performed at least 24 h after
preparation. During this lap of time, the samples were let at
rest at room temperature. All the prepared polymer solutions
were clear, showing that no macroscopic aggregation was
present in the solutions.
Rheometry. The linear and nonlinear rheological properties of the polymer solutions were studied using a stresscontrolled Rheometric Scientific SR 200, equipped with either
a cone and plate geometry (diameter ) 25 mm, cone angle )
5.7°, truncation ) 56 µm) or a Couette geometry (gap ) 1.1
mm) depending on the viscosity of the solutions. After loading,
each sample is let at rest for 5 min before measurements to
remove the mechanical history. Viscosity measurements were
taken in an “equilibrium” state of the samples under shear,
based on the condition that the time evolution of the shear
rate was smaller than 1%/s. If this condition was not respected,
a limited time of 100 s was chosen to avoid too long measurement. The temperature is fixed at 25 ( 0.1 °C, and the samples
were enclosed in a small volume to prevent them from
prospective water evaporation.
Results
Nonlinear Behavior. Viscosity measurements have
been carried out in a concentration range from 0.16 to
6 wt %. Figure 1 shows the apparent viscosity vs shear
stress for some polymer solutions studied in this con-
Figure 2. Zero-shear viscosity as a function of polymer
concentration.
centration range. For each concentration, a cycle of
increasing shear stress (full symbols) and decreasing
shear stress (open symbols) has been applied. Three
concentration regimes can be defined according to
viscous behavior observed. Below a concentration Cg of
2.5 wt %, corresponding to the semidilute regime,
polymer solutions exhibit a Newtonian response in the
shear stress range studied. Above Cg, the steady shear
viscosity profile depends on the way the shear stress is
applied and also on the concentration range. With an
increasing shear stress, first a Newtonian plateau is
observed, followed by a shear thickening effect, and then
the viscosity decreases at elevated shear stress. The
shear thinning effect is very pronounced above C′ ) 4.5
wt %. This concentration C′ marks the beginning of a
third concentration regime. The change in the flow
behavior around C′ is more obvious when the applied
shear stress decreases. In fact, the viscosity increases
continuously for a concentration below this critical
concentration and overtakes the initial value at low
shear stress while the viscosity reaches the initial
Newtonian value at low shear stress for concentrations
above C′.
Newtonian viscosities obtained from Figure 1 by
increasing the shear stress are plotted in Figure 2 as a
function of polymer concentration. The three concentration regimes discussed previously are clearly identifiable
in this representation. A semidilute regime, an intermediate regime, and a more concentrated regime are
described by a power law dependence of the Newtonian
Figure 3. Viscosity as a function of shear stress for a 4 wt %
polymer solution. Cycle of increasing and decreasing shear
stress in the Newtonian region (a) and up to the shear
thickenig region (b). The cycle has been described using the
notation of Figure 1.
Figure 4. Thixotropic flow curve of 4 wt % polymer solution.
Square symbols and circle symbols correspond respectively to
the first and the second flow cycle. The cycle has been
described using the notation of Figure 1.
viscosity with an exponent of about 0.58, 19, and 5.4,
respectively. Only the exponent of 0.58, observed in the
semidilute regime, is in good agreement with the scaling
theory of semidilute unentangled polyelectrolyte solutions.18 However, from a qualitative point of view, a
similar dependence of Newtonian viscosity with concentration has been observed for hydrophobically associative polymer solutions forming micellar gel.19-21
To have some physical insight into the microstructure
in the intermediate regime, a complementary study has
been carried out at Cp ) 4 wt %. Figure 3 shows an
increasing/decreasing stress sweep test in the linear
regime (a) corresponding to the Newtonian behavior and
up to the shear thickening regime (b). This figure clearly
shows that the enhancement of viscosity with decreasing stress is only in the nonlinear regime. Two consecutive cycles of increasing and decreasing shear stress
have been applied in an extended shear stress range.
Figure 4 shows the viscosity of the solution as a function
of the shear stress. After the first cycle, the viscosity at
low stress is 3 times higher than that at the initial state.
Preshearing has modified irreversibly the initial structure. In the second cycle, the shear thickening effect has
been vanished, and the viscosity profile obtained by
increasing the shear stress is simply characterized by
an expanded Newtonian plateau followed by a shearthinning effect. When the applied shear stress de-
Rheological Properties of Physical Gel 3901
Figure 5. Storage modulus G′ and loss modulus G′′ as a
function of shear strain for 4 and 6 wt % polymer solutions.
creases, the flow curve is similar to that observed during
the first cycle. These original results show that the
thickening effect depends on the mechanical history
imposed on the structure.
Linear Behavior. The strain dependence of the
storage modulus and the loss modulus has been first
measured in order to determine the linear viscoelastic
regime. Figure 5 represents the storage modulus and
the loss modulus as a function of strain for 4 and 6 wt
% polymer solutions measured at the frequency of 0.5
Hz. As the strain increases, both moduli exhibit a
plateau value G′0 and G′′0 until a critical strain γc of
about 50%, above of which moduli increase, reach a
maximum value, and decrease. It has to be noted that
the strain hardening of the storage modulus appears
at a lower strain than that of the loss modulus. While
such a peak in the loss modulus profile has been
observed for associative polymer solutions,8,9 a peak in
the storage modulus is in great contrast to those
generally observed, characterized by a pronounced
decrease of G′ with increasing strain beyond the linear
viscoelastic regime. According to our knowledge, only
HASE solutions present a similar strain dependence of
G′ and G′′ moduli, which however do not exhibit a shear
thickening effect in the shear viscosity profile.10-13 For
these systems, the strain hardening increases with
increasing the length of the hydrophobic groups, i.e., the
strength of the associative junction. Moreover, these
peaks are only observed at high frequency for which the
polymer network does not have sufficient time to relax
within the time of one oscillation cycle. Dynamic measurements have been extended in the concentration
range from 3.5 to 6 wt %.
Our interest focuses on the unusual strain dependence of the G′ modulus. Figure 6 represents the storage
modulus as a function of strain at different polymer
concentrations. All the polymer solutions studied in this
concentration range exhibit a peak in the storage
modulus. Let us try to compare the shear thickening
effect (shear-induced viscosity enhancement) and the
strain hardening of the G′ modulus. At a concentration
of 4 wt %, the shear thickening effect arises at a shear
rate of about 0.2 s-1. In dynamic measurements, the
product γc f is the shear rate reached at the critical shear
strain amplitude, with f the frequency. This shear rate,
of about 0.25 s-1, is in the same order of magnitude than
that obtained in steady shear flow. So, from a phenomenological point of view, the peak in the storage modulus
has some similarity with the shear-thickening behavior
3902 Bossard et al.
Figure 6. Storage modulus G′ as a function of shear strain
for 3.5, 4, 4.5, 5, 5.5, and 6 wt % polymer solutions.
Figure 8. Reduced storage modulus G′/G′0 as a function of
shear strain for 3.5 (b), 4 (0), 4.5 (1), 5 (]), 5.5 (9), and 6 (4)
wt % polymer solutions. Inset: intensity of the reduced storage
modulus G′max/G′0 as a function of polymer concentration.
Figure 7. Plateau values of the storage modulus G′0 as a
function of polymer concentration.
observed in the weakly nonlinear part of the flow curve.
Besides, it occurs that both viscoelastic moduli are
significantly enhanced by polymer addition.
This concentration dependence is illustrated in Figure
7, showing the plateau storage modulus as a function
of polymer concentration. It is noted that the term
“plateau” is used in this work to qualify the linear
viscoelastic modulus obtained at a fixed frequency (0.5
Hz). The concentration dependence of the plateau storage modulus has been described using a power law with
an exponent of 9.3 below C′ and 4.6 above C′. The
Newtonian viscosity and the storage moduli are material properties, which exhibit the same qualitative dependence as a function of concentration. In the third
regime, these material properties increase more slowly
than in the intermediate regime. It should be noted that
the exponents of the power laws in both material
properties beyond the percolation threshold are much
higher than those observed in other water-soluble
polymeric thickeners, approaching the theoretical predictions of Semenov et al.20
Let us now look at the qualitative behavior of the
storage modulus. Figure 8 shows the reduced storage
modulus as a function of shear strain for concentration
between 3.5 and 6 wt %. The arising of the strain
hardening appears at the strain amplitude of about 50%
for the storage modulus and does not depend on polymer
concentration. This figure points out a significative
dependence of the peak intensity of the reduced storage
modulus, noted G′max/G′0, with polymer concentration.
The inset shows the peak intensity of the reduced
storage modulus as a function of polymer concentration.
Figure 9. Storage modulus G′ and loss modulus G′′ as a
function of frequency for 3.5, 4.5, and 6 wt % polymer solutions.
The peak intensity of the storage modulus decreases
significantly with increasing concentration up to about
5 wt % and then levels off.
The variation of dynamic moduli with frequency has
been measured in the linear viscoelastic range. Figure
9 shows the dynamic moduli as a function of frequency
for 3.5, 4.5, and 6 wt % polymer concentrations. Two
viscoelastic behaviors can be distinguished in accordance with concentration regimes: in the intermediate
regime both modules depend on frequency, showing that
the solutions behave like a viscoelastic liquid. On the
contrary, in the third regime dynamic moduli are
practically independent of frequency, at least in the
range of frequency studied, showing that concentrated
solutions behave like an elastic gel.
In all concentrations studied it was not possible to
observe the terminal zone of the relaxation spectrum.
Therefore, we were not able to apply any model to
determine the relaxation behavior of the system. However, a rough estimation of the longest relaxation time,
τ, could be given by the intersection of the storage and
loss modulus curves. In the third concentration regime
τ is of the order of 300 s, and it seems to be independent
of polymer concentration.
Discussion
In the dilute regime (C < Cg ) 2.5 wt %), the
polyampholyte is expected to form aggregates due to the
electrostatic interactions between oppositely charged
blocks. The elucidation of the microstructure of the
associated macromolecules below Cg will be the subject
of a forthcoming publication. Although the structure of
these aggregates has not been yet investigated, we could
assume two possible mode of association.
Negatively charged PAA units at both ends of the
polymer form polyelectrolyte complexes with the neighboring positively charged P2VP units within a single
copolymer molecule. Because of the hydrophobic nature
of these polyelectrolyte complexes, the polymer is transformed to a hydrophobically end-capped polyelectrolyte,
which could form “flowerlike” micelles with loops of
extra P2VP charged chains in the corona. By increasing
concentration, a network of bridging micelles should be
formed.22 This scenario requires the existence of high
negative charge density in the PAA end-blocks, which
it is not favored at pH 3.4. Moreover, the rheological
behavior of such a system should resemble to that of
telechelic polyelectrolytes,7,8 whereas it differs in many
respects.
Another more plausible association mechanism could
be suggested. Some negative charges on PAA blocks
interact with the positive charges located along the
major P2VP protonated middle chain imposing intra
and/or intermolecular associations (Figure 10a). This
leads to the formation of open loose aggregates without
discernible hydrophobic domains since the polyelectrolyte complexes are short due to the limiting number of
negative charges in the PAA blocks. Above Cg, a
percolation process due to electrostatic associations
between oppositely charged groups leads to the formation of a loose network (Figure 10b). Beyond Cg, the
rheological properties of this network are highly enhanced by polymer addition while they increase to a
lesser extend above C′. Let us discuss now the nature
of this network. The value of γc, close to 50%, is rather
important in comparison to that measured for telechelic
polyelectrolyte solutions (γc ∼ 1%).8 The short value of
the linear viscoelastic range has been attributed to a
stretched conformation of the polyelectrolyte arising
from intramolecular electrostatic repulsions. The high
value of γc obtained with PAA135-P2VP628-PAA135
tends to prove that the polyampholyte is not fully
stretched. This result is in good agreement with recent
AFM observations of P2VP polymer at pH 3.4.23 A
second point to be mentioned is some similarities
observed in the rheological behavior of HASE, i.e., strain
hardening at the upper limit of the linear viscoelastic
range, local shear thickening followed by a shear thinning beyond a critical shear stress, and a sharp concentration dependence of the Newtonian viscosity.
We present a possible microstructural interpretation
of the rheological behavior. The character of the polymer
under investigation is hydrophilic (lack of hydrophobic
stickers), and this is a fundamental difference from the
amphiphilic character of the associative polymers studied so far. The transient network is mainly composed
of elastically active chains arisen from intermolecular
electrostatic interactions between negatively charged
groups located in the end PAA blocks and positively
charged protonated pyridine groups located in the
middle long P2VP block.16 Intramolecular associations
are likely to occur since at pH 3.4 the long P2VP chains
do not adopt exclusively a stretched conformation as we
mentioned above. Moreover, a number of PAA block
ends may stay unassociated (dangling ends) as they are
water-soluble, contrary to what occurs in hydrophobi-
Rheological Properties of Physical Gel 3903
Figure 10. Schematic representation of the possible molecular microstructure at rest in (a) the dilute regime, (b) the
intermediate regime, and (c) the more concentrated regime.
cally associative polymers. The last two cases lead to a
number of elastically inactive chains (dead branches for
the network) that nevertheless can be considered as
potentially active in the rheological behavior of the
solutions (Figure 10b). It is admitted that the storage
modulus reflects the number density of the elastically
active chains. As the strain amplitude increases beyond
γc, the dangling ends are forced to form new electrostatic
junctions. Moreover, competition between dissociations
and associations promotes extra intermolecular associations due to an intra- to intermolecular transition. Both
processes lead to an increase of G′ modulus and shear
viscosity. The proposed microstructural analysis is
corroborated by the concentration dependence of the
peak intensity of the reduced dynamic moduli: by
increasing polymer concentration, the decrease of the
peak intensity of the storage modulus may be attributed
to the progressive reduction of the number of extra
intramolecular associations. Finally, as the strain am-
3904 Bossard et al.
plitude increases further, the dissociation rate is more
important than the association rate; the network collapses leading to a drop in the dynamic moduli.
Using the microstructural organization proposed here,
the influence of the mechanical history on the viscous
behavior could be interpreted as follows: At the intermediate shear stress corresponding to the shearthickening effect, the shear flow forces both intramolecular associated and “dangling” ends to form new
elastically active intermolecular bridges which harden
the transient network. At high shear stress, the gradually fragmentation of the network leads to a decrease
of the viscosity. As the shear stress decreases from high
stresses, intermolecular junctions are gradually created
leading, at low stress, to a new transient network
characterized by a higher number of bridges than in the
initial structure. The relaxation time, i.e., the time
needed to achieve this more structured organization, of
about 300 s, is rather long in comparison to the
experiment time and could explain the lower viscosity
observed in this rebuilding procedure. When this new
structure is submitted to an increasing stress, there is
only a competition between association and dissociation
of efficient associative junctions from the network. The
shear flow does not increase the number of elastically
active bridges, and therefore no shear-thickening effect
is observed.
In summary, the rheological results suggest that
below C′ a loose network is formed, allowing pronounced
thickening effects by concentration enhancement and/
or shearing, while above C′, a complete 3-dimensional
network almost free from “dangling ends” is achieved
as is illustrated in Figure 10.
Concluding Remarks
The rheological properties of a physical gel formed by
an asymmetric triblock polyampholyte of the type
PAA135-P2VP628-PAA135 in salt-free aqueous solutions
have been presented. To the best of our knowledge, it
is shown for the first time that a weak polyampholyte
with asymmetric triblock copolymer architecture can
form an infinite three-dimensional reversible network
in a certain pH. It is also the first example of a double
hydrophilic block copolymer (lack of hydrophobic stickers) that behaves as a strong thickener (6 orders of
magnitude higher viscosity at 4 wt % polymer concentration).
Particular attention has been paid to the association
mechanism and the structure of the transient network
formed above Cg ) 2.5 wt %. The molecular dynamics
of the network exhibits some particular behavior that
differs from other associative polymer solutions (peak
in both G′ and G′′ moduli in intermediate strain
amplitude, prolonged shear thickening in moderated
shear stress). The whole set of rheological data support
the coexistence in the network of elastically active and
inactive polymer chains named “dead branches” attributed to dangling ends and intramolecular association. Shear induces structural rearrangements by promoting the intra- to intermolecular associations and
forcing the dangling ends to join the mechanically active
network. The rheological properties of the system are
strongly concentration dependent. Two concentration
regimes can be identified above the percolation concentration Cg with characteristic flow behavior. The transition concentration C′ between these regimes has been
interpreted in term of mechanically efficient connectivity of the network.
(i) For concentrations below C′, the network contains
many dangling ends and intramolecular associations
(elastically inactive branches). The rheological properties are strongly improved by polymer addition, which
enhances the network connectivity. The exponents of
η0 and G0 power laws are the highest ever observed in
associative polymers.
(ii) Above C′, a complete 3-dimensional network is
achieved for which polymer addition improves to a lesser
extent its rheological properties.
Acknowledgment. This work has been performed
with the financial support of the European Community
through Contract HPMD-CT2000-00054-02. The contribution of Vasiliki Sfika was performed in the framework of the Operational Program for Education and
Initial Vocational Training on Polymer Science and
Technology of the University of Patras, through the
Ministry of Education and Religious Affairs in Greece.
References and Notes
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(17) Nisato, G.; Candau, S. J. In Polymer Gels and Networks;
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MA0353890
Thickening effect in soluble hydrogen-bonding
interpolymer complexes. Influence of pH
and molecular parameters
F. Bossarda)
Institute of Chemical Engineering and High Temperature Chemical Processes,
ICE/HT-FORTH, P.O. Box 1414, 26504 Patras, Greece
M. Sotiropoulou and G. Staikos
Department of Chemical Engineering, University of Patras, 26500 Patras, Greece
(Received 14 November 2003; final revision received 7 April 2004)
Synopsis
Linear and nonlinear viscoelastic properties of poly共acrylic acid兲共PAA兲 and poly共acrylic
acid-co-2-acrylamido-2-methylpropane
sulfonic
acid兲-graft-poly共N,N-dimethylacrylamide兲
共P共AA-co-AMPSA兲-g-PDMAM兲 mixtures have been investigated as a function of pH, the
PDMAM content of the graft copolymer and the molecular weight of PAA. At pH ⬍ 3.75, strong
hydrogen-bonding interpolymer complexation between PAA and PDMAM side chains in semidilute
solution leads to the formation of a transient network, as the considerable increase in viscosity
indicates. The sol/gel transition observed at pH ⫽ 2.0 by increasing the graft copolymer
composition in PDMAM is explained by a substantial increase in the number of the junctions
共stickers兲 resulting from the PDMAAM/PAA hydrogen bonding complexation. Moreover, the
thickening effect observed is further strengthened by increasing the molecular weight of PAA, due
to the interconnection of more copolymer chains. © 2004 The Society of Rheology.
关DOI: 10.1122/1.1763941兴
I. INTRODUCTION
The thickening property of water-soluble polymer solutions is required in many industrial applications, for example in pharmaceutics, coatings industry, food, and oil recovery 关Glass 共1989, 1991兲; Shalaby et al. 共1991兲; Goddard and Gruber 共1999兲兴. A pronounced thickening effect occurs when polymer chains self-associate in a transient
network through temporary junctions between functional groups. Traditionally, these
functional groups are long hydrocarbon or fluorocarbon alkyl chains 关Schulz et al.
共1987兲; McCormick et al. 共1988兲; Wang et al. 共1988兲; Hill et al. 共1993兲; Petit et al.
共1996兲; Volpert et al. 共1998兲; Ma and Cooper 共2001兲; Tsitsilianis and Iliopoulos 共2002兲兴
or thermo-sensitive chains, such as poly共ethylene oxide-co-propylene oxide兲关Vos et al.
共1994兲兴, poly共ethylene oxide兲关Hourdet et al. 共1994兲兴 or poly共N-isopropylacrylamide兲
共PNIPAM兲关Bokias et al. 共1997兲; Durand and Hourdet 共1999兲; Bokias et al. 共2001兲兴. The
rheological properties of such associative polymers have been thoroughly studied during
a兲
Author to whom correspondence should be addressed; electronic mail: [email protected]
J. Rheol. 48共4兲, 927-936 July/August 共2004兲
0148-6055/2004/48共4兲/927/10/$25.00
927
928
BOSSARD, SOTIROPOULOU, AND STAIKOS
the past few years 关Wang et al. 共1991兲; English et al. 共1997, 1999兲; Regalado et al.
共1999兲; Aubry et al. 共2002, 2003兲兴.
Hydrophobically modified copolymers 共mostly acrylic acid based兲 have been studied
and it has been shown that pH plays an important role in the viscosity enhancement of
these systems, even if the hydrophobic effect is the driving force for the formation of
hydrophobic domains cross-linking the different polymer chains 关Smith and McCormick
共2001兲; Li and Kwak 共2002兲兴. A pronounced pH-dependent thickening effect has been
recently observed for a tri-block copolymer composed of a central block of poly共2vinylpyridine兲 end-capped by two polyacrylic acid 共PAA兲 blocks 关Sfika and Tsitsilianis
共2003兲兴.
However, a thickening effect induced by pure hydrogen-bonding interactions has been
observed only in mixtures of high molecular weight neutral polybases, such as poly共ethylene oxide兲 or polyvinylpyrrolidone, with charged copolymers of acrylic acid 共AA兲
关Iliopoulos and Audebert 共1985兲; Iliopoulos et al. 共1988兲; Iliopoulos and Audebert
共1991兲兴. If these copolymers are just of partially neutralized PAA the enhancement of
viscosity appears in a narrow pH range, as at pH higher than 4 –5 hydrogen bonding
complexation is not possible, while at pH lower than 3–3.5 compact hydrogen bonding
interpolymer complexes precipitate 关Ikawa et al. 共1975兲; Eustace et al. 共1988兲; Usaitis
et al. 共1997兲兴. In the case of AA copolymers with strongly ionized monomers, such as
2-acrylamido-2-methylpropane sulfonic acid 共AMPSA兲, the copolymerization degree
must be high enough to avoid precipitation at low pH but not too high to prevent
complexation 关Iliopoulos and Audebert 共1991兲兴. In these systems a delicate compromise
has to be achieved between the necessary complexable carboxylic units and the adequate
anionic hydrophilic groups on the same chain, so that adequate complexation and swelling would simultaneously occur leading to gel formation. Recently, we have reported on
the thickening behavior of a polymer mixture of the graft copolymer P共AA-co-AMPSA兲g-PDMAM with PAA in semidilute aqueous solution, which showed a considerable
thickening effect by decreasing pH, due to the formation of a stable transient network
through hydrogen bonding complexation 关Sotiropoulou et al. 共2003兲兴. In this system, the
poly共N,N-dimethylacrylamide兲共PDMAM兲 side chains and the PAA homopolymer constitute the hydrogen bonding complexable agents 关Wang and Morawetz 共1989兲兴, while the
negatively charged AMPSA units, present along the graft copolymer backbone, provide a
sufficient hydrophilicity, so that precipitation of the interpolymer complexes formed at
low pH is avoided 关Sotiropoulou et al. 共2003兲兴. 共The low percentage of AA units have
been introduced in the graft copolymer backbone just to provide the necessary functional
carboxylic groups for grafting of the PDMAM side chains.兲
In this work, the pH dependency of the viscoelastic behavior of this novel interpolymer complex is thoroughly studied. Moreover, the influence of the graft copolymer composition in PDMAM side chains and the molecular weight of PAA on the rheological
properties of this system have also been investigated.
II. MATERIALS AND EXPERIMENTS
A. Materials
The two samples of PAA used, PAA90 共Aldrich兲 and PAA450 共Polysciences兲, with an
average molecular weight of 9.0⫻105 g/mol and 4.5⫻105 g/mol, respectively, were purified with a Pellicon tangential flow filtration system 共Millipore兲 equipped with an ultrafiltration membrane 共Millipore, cutoff 100.000 g/mol兲 and freeze-dried.
The monomers AA, AMPSA and N,N-dimethylacrylamide 共DMAM兲 were purchased
from Aldrich. Ammonium persulfate 共APS, Serva兲, potassium metabisulphite 共KBS, Al-
HYDROGEN-BONDING INTERPOLYMER COMPLEXES
929
drich兲,
2-aminoethanethiol
hydrochloride
共AET,
Aldrich兲
and
1-共3共dimethylamino兲propyl兲-3-ethyl-carbodiimide hydrochloride 共EDC, Aldrich兲 were used
for the synthesis of the graft copolymers.
For the preparation of the buffer solutions citric acid 共CA兲 and Na2 HPO4 from Merck
were used.
Water was purified using a Seralpur Pro 90C 共Germany兲 apparatus combined with a
USF Elga laboratory unit.
B. Polymer synthesis and characterization
Amine-terminated PDMAM was synthesized by free radical polymerization of
DMAM in water at 30 °C for 6 h using the redox couple APS and AET as initiator and
chain transfer agent, respectively. The polymer was purified by dialysis against water
through a membrane with a molecular weight cutoff 12.000 g/mol 共Sigma兲 and finally
obtained by freeze-drying. Its number average molecular weight was determined by end
group titration with NaOH after neutralization with an excess of HCl, using a Metrohm
automatic titrator, model 751 GPD Titrino, and found equal to 17.000 g/mol.
A copolymer of AA and AMPSA, P共AA-co-AMPSA兲 was prepared by free radical
copolymerization of the two monomers in water, after partial neutralization 共90% mole兲
with NaOH at pH ⬃ 6 – 7, at 30 °C for 6 h using the redox couple APS/KBS. The
product obtained was then fully neutralized (pH ⫽ 11) with an excess of NaOH, purified
by ultrafiltration with the above Pellicon system and received in its sodium salt form after
freeze-drying. Its composition, determined by elemental analysis 共Carlo-Erba CHNS-O
elemental analyzer EA 1108兲, was found 20% in AA units. Its weight average molecular
weight, M w ⫽ 1.5⫻105 , was determined by static light scattering in 0.1 M NaCl with a
spectrogoniometer, model SEM RD 共Sematech, France兲 equipped with a He–Ne laser
共633 nm兲. The required refractive index increment, dn/dc ⫽ 0.153, was determined by a
Chromatix KMX-16 He–Ne laser differential refractometer.
The graft copolymers, P共AA-co-AMPSA兲-g-PDMAM22, P共AA-co-AMPSA兲-gPDMAM42, and P共AA-co-AMPSA兲-g-PDMAM60 were synthesized by a coupling reaction between P共AA-co-AMPSA兲 and amine-terminated PDMAM. The two polymers
were dissolved in water in a 5% solution. Then, a fivefold excess of the coupling agent,
EDC, was added and the solution was stirred for 6 h at room temperature. Addition of
EDC was repeated for a second time. The products were purified with the Pellicon system
as above and freeze dried. Their percentage weight composition in PDMAM was determined by means of carbon, nitrogen, and sulfur elemental analysis and on the basis of the
chemical type of the different repeating units, and it is expressed by the number at the end
of their name, with an error estimated at ⫾2.
A schematic description of the graft copolymers and the chemical structure of their
monomers are proposed in Fig. 1.
C. Solutions preparation
PAA and P共AA-co-AMPSA兲-g-PDMAMx 共x is the percentage weight composition in
PDMAM of the graft copolymer兲 6 wt% solutions were initially prepared in 0.15 M citric
acid/phosphate buffers. Then, they were mixed at a 1:1 ratio. Total polymer concentration, 6 wt%, was much higher than the overlap concentration of the anionic backbone of
the graft copolymers, estimated at 0.9 wt% from its intrinsic viscosity, 关 ␩ 兴
⫽ 115 cm3 /g, in a 0.15 M citric acid solution. The mixtures were stirred for 24 h, before
rheological measurements. Their pH was measured with a Metrohm model 713 pH Meter
equipped with a Metrohm combined pH glass needle electrode.
930
FIG. 1. Schematic picture of the P共AA-co-AMPSA兲-g-PDMAM graft copolymer and the chemical structure of
its monomers.
D. Rheometry
The linear and nonlinear rheological measurements were carried out using a Rheometrics SR 200 controlled-stress rheometer, equipped with a cone and plate geometry
(diameter ⫽ 25 mm, cone angle ⫽ 5.7°, truncation ⫽ 56 ␮ m). After loading, each
sample was kept at rest for 5 min before measurement to remove the mechanical history.
Viscosity measurements were taken in steady shear flow state, based on the condition that
the time evolution of the shear rate was smaller than 1%/s. If this condition was not
respected, a limited time of 100 s was chosen to avoid too long measurements. The
temperature was fixed at 25⫾0.1 °C and the samples were enclosed in a small volume to
prevent solvent evaporation.
III. RESULTS AND DISCUSSION
The pH dependence, the influence of the graft copolymer weight composition in
PDMAM, and the influence of the PAA molecular weight on the rheological behavior of
the mixtures of the graft copolymers P(AA-co-AMPSA)-g-PDMAMx with PAA in
semidilute solutions have been studied.
A. p H dependence
Viscoelastic measurements, carried out for the P共AA-co-AMPSA兲-g-PDMAM60/
PAA450 polymer mixture at pH values 2.0, 3.4, and 3.8, are shown in Fig. 2, presenting
typical frequency sweeps and revealing the strongly pH-dependent thickening behavior
of the polymer mixture studied. At pH 3.8 the mixture behaves like a viscoelastic liquid,
as the pulsation dependencies of G⬘ and G⬙ at low frequencies are proportional to ␻ 2 and
␻ 1 , respectively. For dense molecular systems, long time dynamics is governed by reptation 关De Gennes 共1979兲兴. The reptation time, roughly corresponding to the G⬘ -G⬙
crossover, is lower than 0.06 s. At pH 3.4, the viscoelastic behavior of the mixture is
similar to that at pH 3.8 but G⬘ and G⬙ moduli are greatly enhanced and the reptation
931
FIG. 2. Linear viscoelastic behavior of P共AA-co-AMPSA兲-g-PDMAM60/PAA 450 mixtures in buffer solutions
of pH 2.0, 3.4, and 3.8. The total polymer concentration is 6% w/w and the weight ratio of the two polymers
in the mixture is 1:1, the same as for all polymer mixtures studied in this work. Storage modulus, G⬘ , 䊏; loss
modulus, G⬙ , 䊐.
time attains 0.6 s, showing a pronounced slowing down of the molecular dynamic. At pH
2.0 the mixture behaves like a gel, as G⬘ modulus is higher than G⬙ modulus and both are
practically pulsation independent. This sol/gel transition ensues from the graft copolymer
architecture, combining the proton acceptor ability of the PDMAM side chains, which
form strong hydrogen-bonding interpolymer complexes with PAA 关Wang and Morawetz
共1989兲; Shibanuma et al. 共2000兲兴, and the highly hydrophilic character of the copolymer
backbone, resulting from its high percentage composition 共80%兲 in the strongly anionic
AMPSA units. Here 3.75 is a critical pH value because hydrogen bonding interpolymer
complexation between PAA and PDMAM is prevented at pH ⬎ 3.75 关Sotiropoulou
et al. 共2003兲兴. By decreasing pH, hydrogen-bonding complexation between PAA and the
PDMAM side chains is progressively strengthened, resulting in the formation of stickers
along the graft copolymer anionic backbone, which hinder the molecular dynamic and
enhance the viscoelastic properties 关Rubinstein and Semenov 共2001兲; Semenov and Rubinstein 共2002兲兴. At pH ⫽ 2.0, the PDMAM/PAA junctions are much more strengthened,
due to the strong hydrogen bonding complexation, leading to precipitation in the case of
the two pure homopolymers 关Sotiropoulou et al. 共2003兲兴. Nevertheless, our interpolymer
complex remains soluble, due to the negatively charged AMPSA units that are the major
constituents of the graft copolymer backbone. Consequently, a very interesting complex
system is formed, comprised by insoluble hydrogen-bonding interpolymer complexes of
PAA with the PDMAM side chains of the graft copolymer, functioning as stickers and
binding the hydrophilic, well-extended negatively charged graft copolymer backbones,
resulting in a gel-like behavior.
B. Influence of the graft copolymer weight composition in PDMAM
Figure 3 shows the steady-state viscosity versus shear stress of the
P共AA-co-AMPSA兲-g-PDMAM22/PAA450, P共AA-co-AMPSA兲-g-PDMAM42/PAA450,
and P共AA-co-AMPSA兲-g-PDMAM60/PAA450 polymer mixtures at pH ⫽ 2.0.
For each solution, a cycle of increasing 共full symbols兲 and decreasing shear
stress 共open symbols兲 has been applied. The viscous profiles observed for the
932
FIG. 3. Viscosity versus shear stress for the P共AA-co-AMPSA兲-g-PDMAM22/PAA 450 共䊉, 䊊兲; P共AA-coAMPSA兲-g-PDMAM42/PAA 450 共䉱, 䉭兲 and P共AA-co-AMPSA兲-g-PDMAM60/PAA 450 共䊏, 䊐兲 mixtures at
pH 2.0. Full symbols correspond to data obtained by increasing stress and open symbols to that obtained by
decreasing stress.
P(AA-co-AMPSA)-g-PDMAM60 based mixture form a hysteresis loop, characteristic
of thixotropic materials. When the shear stress increases, the polymer solution presents a
high Newtonian viscosity, followed by a discontinuous shear thinning. In the decreasing
stress mode, viscosity increases again, but with values considerably lower than those in
the increasing stress run, a behavior representative of structured systems. On the contrary,
the viscous behavior of the polymer mixtures of P(AA-co-AMPSA)-g-PDMAM22 and
P(AA-co-AMPSA)-g-PDMAM42 with PAA450 exhibits a well-defined Newtonian plateau at low shear stress, followed by a smooth shear thinning with no hysteresis loop.
Such a rheological behavior is observed for polymer solutions with weak intermolecular
interactions, as for neutral polymer solutions in the semi-dilute regime, where entanglements occur, or in weakly associative polymer solutions 关Aubry and Moan 共1996兲兴. These
results reveal that the mixture of the P(AA-co-AMPSA)-g-PDMAM60 graft copolymer
with PAA450 at pH ⫽ 2.0 is structured, while the corresponding mixtures of
P(AA-co-AMPSA)-g-PDMAM22 and P(AA-co-AMPSA)-g-PDMAM42 appear simply as dense macromolecular systems.
The dynamics of the same systems have been also investigated by linear viscoelastic
measurements. Figure 4 shows their storage, G⬘ , and loss modulus, G⬙ , at pH ⫽ 2.0 as
a function of pulsation. As the weight composition of the graft copolymer in PDMAM
increases, the rheological behavior changes drastically from a viscous liquid (G⬘
⬍ G⬙ and G⬘ ⬃ ␻ 2 and G⬙ ⬃ ␻ 1 ) to a weak elastic solid (G⬘ ⬎ G⬙ and nearly
independent of pulsation in the window studied兲, accompanied by a pronounced increase
in the values of the moduli. The increase of the graft copolymer composition in PMDMAM corresponds to an increase of the number of PDMAM side chains grafted onto the
anionic graft copolymer backbone. Therefore, the sol/gel transition, observed with increasing the weight composition of the graft copolymer in PDMAM, results from an
increase of the number of stickers formed between the P(AA-co-AMPSA)-g-PDMAMx
chains and the PAA chains. It is obvious that a critical number of stickers is needed for
the formation of a transient network.
933
FIG. 4. Linear viscoelastic behavior of the P共AA-co-AMPSA兲-g-PDMAM22/PAA 450 共䊉, 䊊兲; P共AA-coAMPSA兲-g-PDMAM42/PAA 450 共䉱, 䉭兲 and P共AA-co-AMPSA兲-g-PDMAM60/PAA 450 共䊏, 䊐兲 mixtures at
pH 2.0. Full symbols correspond to the storage modulus, G⬘ , and open symbols to the loss modulus, G⬙ .
C. Influence of the molecular weight of PAA
The influence of the molecular weight of PAA on the rheological properties of the
mixtures has been investigated in the linear and nonlinear regime. Figure 5 shows the
steady shear viscosity as a function of the shear stress for two solutions obtained by
mixing the P(AA-co-AMPSA)-g-PDMAM42 graft copolymer with PAA90 and
PAA450 respectively, at pH 2.0. Beyond a linear response, both solutions present a
shear-thinning effect. The increase in Newtonian viscosity with increasing the PAA molecular weight reflects the ability of the longer PAA chains to connect more PDMAM side
chains and as a result to interconnect more graft copolymer chains. This is confirmed by
the linear viscoelastic behavior of these two mixtures, plotted in Fig. 6. The dynamic
moduli are not significantly increased with increasing the molecular weight of PAA,
FIG. 5. Viscosity of the P共AA-co-AMPSA兲-g-PDMAM42/PAA90 共䊉, 䊊兲 and P共AA-co-AMPSA兲-gPDMAM42/PAA 450 共䊏, 䊐兲 mixtures as a function of shear stress at pH 2.0. The cycle of increasing and
decreasing shear stress has been described using the notation of Fig. 2.
934
FIG. 6. Linear viscoelastic behavior of P共AA-co-AMPSA兲-g-PDMAM42/PAA 90 共䊉, 䊊兲 and P共AA-coAMPSA兲-g-PDMAM42/PAA 450 共䊏, 䊐兲 at pH 2.0.
however the molecular dynamics is substantially decreased, with a reptation time passing
from 0.7 s with PAA90 to 40 s with PAA 450. The density of stickers should be the same,
but a larger number of the graft copolymer chains is interconnected by means of the
higher molecular weight PAA, not favoring the reptation dynamics. An overview is
presented in Fig. 7, where the Newtonian viscosity of the graft copolymer mixtures with
PAA90 and PAA450 is plotted as a function of the graft copolymer weight composition
in PDMAM. The Newtonian viscosity increases with increasing the graft copolymer
weight composition in PDMAM, at almost the same rate for the mixtures with PAA90
and PAA450.
IV. CONCLUDING REMARKS
The rheological behavior of mixtures of the P(AA-co-AMPSA)-g-PDMAM graft
copolymer with PAA in semidilute aqueous solution at different pH values, as a function
FIG.
7.
Newtonian
viscosity
of
P(AA-co-AMPSA)-g-PDMAMx/PAA
90
共䊉兲
and
P(AA-co-AMPSA)-g-PDMAMx/PAA 450 共䊏兲 mixtures as a function of the graft copolymer weight composition in PDMAM at pH 2.0.
935
of the graft copolymer composition in PDMAM and at various PAA molecular weights,
has been investigated. A pronounced thickening effect is exhibited:
共a兲
共b兲
共c兲
by decreasing pH, which favors the hydrogen-bonding complexation between the
PDMAM side chains of the graft copolymer and PAA,
by increasing the PDMAM composition of the graft copolymer, which induces an
increase in the number of stickers along the graft copolymer backbone, and
by increasing the molecular weight of PAA, which permits each PAA chain to
interconnect more copolymer chains.
Moreover, a clear transition from a dense liquid to a gel-like solid has been observed
either by decreasing pH in a high PDMAM graft copolymer composition mixture or by
increasing the PDMAM graft copolymer composition at low pH.
To the best of our knowledge, this is the first example of a polymeric system forming
a transient network through pure hydrogen bonding interactions, which are responsible
for a thickening effect, controlled by pH and the molecular parameters of the component
polymers.
ACKNOWLEDGMENT
This work has been performed with the financial support of the European Community,
through Contract No. HPMD-CT2000-00054-02.
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acrylate兲,’’ Polym. Bull. 20, 577–582 共1988兲.
Wang, Y., and H. Morawetz, ‘‘Fluorescence study of the complexation of poly共acrylic acid兲 with poly共N,Ndimethylacrylamide-co-acrylamide兲,’’ Macromolecules 22, 164 –167 共1989兲.
Wang, T. K., I. Iliopoulos, and R. Audebert, ‘‘Aqueous-solution behavior of hydrophobically modified poly共acrylic acid兲,’’ in Water Soluble Polymers. Synthesis, Solution Properties and Applications; S. W. Shalaby,
C. L. McCormick, and G. B. Butler, eds., ACS Symposium Series 467, pp. 218 –231 共1991兲.
Macromolecules 2005, 38, 2883-2888
2883
A Novel Thermothickening Phenomenon Exhibited by a Triblock
Polyampholyte in Aqueous Salt-Free Solutions
Frédéric Bossard,† Constantinos Tsitsilianis,*,†,‡ Spyros N. Yannopoulos,†
Georgios Petekidis,§ and Vasiliki Sfika†,‡
Institute of Chemical Engineering and High-Temperature Chemical ProcessessFoundation for
Research and Technology (ICE/HT-FORTH), P.O. Box 1414, 26504 Patras, Greece;
Department of Chemical Engineering, University of Patras, 26504 Patras, Greece; and
Institute of Electronic Structure and Laser, GR-71110 Heraklion, Crete, Greece
Received December 2, 2004; Revised Manuscript Received February 2, 2005
ABSTRACT: A novel thermothickening phenomenon exhibited by a water-soluble triblock copolymer in
salt-free aqueous solutions has been investigated through rheological measurements and supported by
dynamic light scattering. The copolymer is constituted of a long central poly(2-vinylpyridine) block endcapped by two shorter poly(acrylic acid) blocks, (PAA135-P2VP628-PAA135). At pH 3.4 of interest, the
copolymer behaves as polyampholyte, bearing positively charged protonated 2VP and negatively charged
AA moieties. In aqueous solutions where a physical gel is formed, rheological investigation showed a
pronounced thermothickening behavior upon heating to 35 °C, followed by a weak zero-shear viscosity
decrease. This unexpected temperature dependence has been interpreted by considering a competition
between two antagonistic effects: (i) a remarkable swelling of the macromolecular chain upon heating,
mainly due to the excluded-volume effect of the outer PAA blocks, that favors intermolecular interactions
between oppositely charged blocks responsible for physical gelation and (ii) the thermal motion of molecules
which speeds up the molecular dynamics and tends to weaken the rheological properties. The effect of
the macromolecular swelling prevails at low temperatures while the influence of the thermal motion
increases continually and predominates at high temperatures.
Introduction
Associative polymers, generally used as rheology
modifiers for their thickening properties, are of prime
interest in numerous industrial applications.1,2 Thermoassociative polymers represent a particular class of
self-associating materials since their thickening properties are promoted upon heating. Indeed, such polymers
contain functional groups as poly(ethylene oxide-copropylene oxide),3 poly(ethylene oxide),4 or poly(Nisopropylacrylamide) (PNIPAM)5-8 characterized by a
lower critical solution temperature (LCST). From a
phenomenological point of view, the specific thermal
behavior arises from the progressive formation of a
transient network constituted by hydrophilic backbone
physically cross-linked via hydrophobic interactions of
the thermoresponsive groups at temperatures above
their LCST.5,9,10
Some multiarm star polymers (or colloidal stars)
exhibit thermothickening behavior which differs from
that mentioned above.11-13 These soft materials, bridging the gap between polymer solutions and colloidal
suspensions, display a counterintuitive reversible gelation upon heating. This thermothickening behavior,
which corresponds to a jamming transition, is attributed
to the swelling of dangling polymeric arms when temperature is increased, leading to the gradual interpenetration of the soft spheres. It has been shown experimentally that significant swelling results from the
enhancement of the solvent quality passing from θ
solvent toward good solvent with increasing tempera†
ICE/HT-FORTH.
University of Patras.
§ Institute of Electronic Structure and Laser.
* Corresponding author: Fax +30 2610 997 266; e-mail
[email protected].
‡
ture. The solvent quality can also be improved by
choosing the appropriate solvent for the polymeric arms.
Recently, some of us14a have demonstrated for the first
time that a double hydrophilic water-soluble polymer,
i.e., a triblock copolymer of poly(acrylic acid)-poly(vinylpyridine)-poly(acrylic acid) (PAA-P2VP-PAA),
may be self-organized reversibly into two distinct and
completely different structures (i.e., from a transient
three-dimensional network to compact micelles) by
switching the pH of the aqueous media. This interesting
and novel behavior was attributed to the nature and
the specific architecture of the polymer named asymmetric triblock polyampholyte.14a
In the present work, an unexpected and rather
complex reversible thermothickening phenomenon exhibited by the same polymeric material is presented. At
pH 3.4 of interest in this study, the water-soluble PAAP2VP-PAA copolymer exhibits a polyampholyte character since the central P2VP block is partially protonated and therefore bears positive charges while a
limiting number of negative charges exist in the PAA
blocks due to the acrylic acid dissociation. At T ) 25 °C
and above a critical concentration, a previous work has
shown that this copolymer forms a transient network
through electrostatic interactions between oppositely
charged blocks of different chains14b which probably is
stabilized by hydrogen bonding between the uncharged
moieties. Some “topological defects” in the mechanically
active network, consisting of nonassociated “dangling
ends” and polymeric chains involved in intramolecular
associations, have been considered to account for a
pronounced shear thickening effect. It has been shown
that these “defects” depend on the mechanical history
of the material. For example, a sufficient high shear
stress induces an intra- to interassociation transition
and forces the “dangling ends” to join the mechanically
10.1021/ma047520p CCC: $30.25 © 2005 American Chemical Society
2884
Bossard et al.
active network, leading to a significant zero-shear
viscosity enhancement after a preshearing. On the basis
of this molecular approach, we have extended the
previous study by a thorough investigation of the
temperature dependence of polyampholyte’s rheological
behavior in salt-free water. The unexpected thermoresponsiveness of this polymeric material is compared
with that observed in other systems mentioned above
and discussed in terms of microstructure by using
dynamic light scattering.
Materials. The PAA-P2VP-PAA triblock copolymer was
synthesized by anionic polymerization. Details of the synthesis
are given elsewere.14a The degree of polymerization is 628 for
the central P2VP block and 135 for each PAA end-block, having
total weight-average molecular weight Mw 8.5 × 104 g/mol and
molecular polydispersity Mw/Mn ) 1.11. The copolymer is
water-soluble, and its aqueous solutions give a pH close to 3.4.
The net charge of the polyampholyte is positive (the isoelectric
point was found at pH 5.5) due to the asymmetric architecture
of the macromolecule. Polymer solutions were obtained by
dissolving the proper amount of polymer in distilled water,
and a resting time of 24 h at room temperature was applied
before measurements. All the prepared polymer solutions were
clear, showing that no macroscopic phase separation was
present in the solutions.
Rheometry. Rheological measurements were carried out
using a controlled stress Rheometric Scientific SR 200, equipped
with a cone and plate geometry (diameter ) 25 mm, cone angle
) 5.7°, truncation ) 56 µm). The temperature was controlled
between 10 and 50 °C with an accuracy of (0.1 °C by a water
bath circulator. All rheological measurements were performed
with a 4 wt % polymer solution (c/c* of about 1.6), for which a
physical gel is formed at 25 °C.14b After loading, each sample
was kept at rest for 5 min before measurements to remove
the mechanical history. Viscosity measurements were taken
in a steady-state condition. For this purpose, shear stress
sweep tests were carried out with an equilibration time of 100
s and a time evolution of the shear rate smaller than 1% per
second. To prevent water evaporation, samples were enclosed
in a small solvent trap.
Dynamic Light Scattering. Characterization of the polymeric solution as a function of temperature in the dilute regime
(c ) 0.5 wt %) was achieved via dynamic light scattering. The
normalized intensity time correlation function g(2)(q,t) ) 〈I(q,t)I(q,0)〉/〈I(q)〉2 was measured at various scattering angles and
temperatures spanning a time scale from 10-7 to 103 s. The
measurements were performed with a Nd:YAG laser (ADLAS)
operating at 532 nm with a stabilized power of 100 mW in an
ALV goniometer setup. The polarization conditions of the
incident and scattered radiation were controlled by utilizing
a set of a Glan and Glan-Thomson polarizers (Halle, Berlin)
with an extinction coefficient better than 10-7. The scattered
light was analyzed with a full multiple-τ fast digital correlator
(ALV-5000/E) with 280 channels. The relaxation of concentration fluctuations in the polymer solution due to the Brownian
motion of the polymer was detected at different scattering
wave vectors q ) (4πn/λ) sin(θ/2), where n is the refractive
index of the solvent and θ the scattering angle. The intermediate scattering function, C(q,t), was deduced from g(2)(q,t) by
the Siegert relation15
g(2)(q,t) ) 1 + f *|C(q,t)|2
(1)
where f* is an instrumental factor related to the coherence
area. C(q,t) was analyzed as a weighted sum of independent
contributions
C(q,t) )
∫L(ln τ) exp(-t/τ) d ln τ
(2)
The distribution of relaxation times L(ln τ) was obtained by
the inverse Laplace transformation (ILT) of C(q,t) using the
Figure 1. Dynamic temperature ramp for a 4 wt % solution
monitored at 1 rad/s with a temperature rate of 0.5 °C/min.
Full symbols correspond to data obtained by increasing temperature and open symbols to that obtained by decreasing
temperature.
CONTIN algorithm.16 The apparent hydrodynamic radii of
macromolecules were determined using the Stokes-Einstein
relation
RH )
kBT
6πηD
(3)
where kB is the Boltzmann constant, η is the viscosity of the
solvent, and D is the diffusion coefficient. The latter was
determined by D ) Γ/q2 at the limit q ) 0, where Γ is the decay
rate of C(q,t).
Results and Discussion
Rheology. a. Oscillatory Shear Measurements.
Figure 1 depicts the linear viscoelastic response of the
4 wt % polymer solution subjected to an increasing
temperature ramp (full symbols) followed by a decreasing temperature ramp (open symbols). Measurements
were conducted at a frequency ω ) 1 rad/s in a
temperature range from 10 to 50 °C with a temperature
rate of 0.5 °C/min. It is evident that the polymer solution
exhibits a pronounced reversible thermothickening effect. However, both moduli G′ and G′′ undergo a
hysteresis effect over a complete cycle of temperature
change. This leads to a partial recovery of their values
at the end of the heating/cooling sweep test. This
peculiar temperature dependence observed for a nonthermoassociative polymer could be due to a specific
thermodynamic process characterized by an intrinsic
hysteresis and/or a slow kinetic behavior, at least slower
than the experimental time scale. Further experiments
were performed to characterize thoroughly this unusual
thermoresponsive effect.
We first focus our attention on the behavior of G′ and
G′′ moduli as a function of shear strain. The shear strain
dependence of storage modulus G′ and loss modulus G′′
normalized by their plateau values G′0 and G′′0 are
shown in Figures 2 and 3, respectively, for different
temperatures ranging from 18 to 50 °C. Let us stress
that the term “plateau values” is used in this paper to
qualify the strain-independent (i.e., linear) viscoelastic
moduli determined at a fixed frequency. For all temperatures considered, the solution exhibits a linear
viscoelastic response below a critical shear strain γc ∼
0.3, where both G′ and G′′ moduli are independent of
shear strain. A further increase in the amplitude of the
oscillatory shear strain leads to strain hardening characterized by a pronounced increase in G′ followed by an
appreciable drop at even higher strains.
Triblock Polyampholyte in Aqueous Solutions 2885
Figure 2. Reduced storage modulus G′/G′0 of the 4 wt %
solution as a function of shear strain γ at T ) 18 (9), 22 (O),
27.5 (2), 30 (]), 35 (1), 45 (tilted 4), and 50 °C (tilted 2) and
ω ) 3 rad/s. Inset: (a) Reduced storage modulus G′/G′0 as a
function of shear strain γ. (b) Peak intensity of the reduce
storage modulus G′max/G′0 as a function of temperature. The
full line in inset (a) shows a fit of a second-order power series
in γ2.
Figure 3. Reduced loss modulus G′′/G′′0 of the 4 wt % solution
as a function of shear strain γ at T ) 18 (9), 22 (O), 27.5 (2),
30 (]), 35 (1), 45 (tilted 4), and 50 °C (tilted 2) and ω ) 3
rad/s.
From a molecular point of view, the origin of the
strain hardening evidenced in G′ modulus might be
twofold. As proposed in our previous work,14b the shear
strain at moderated deformations increases the number
of intermolecular interactions between adjacent chains,
which become part of the stress-conducting network,
leading to an enhancement of the storage modulus. This
effect results from a modification of the molecular
conformation which may induce a second effect. The
polymer backbone, adopting a wormlike conformation
at rest, can be stretched under large strains, increasing
its rigidity, and as a result the solution becomes
increasingly stiffer. On the basis of this assumption,
Gisler et al.17 have proposed a model to describe the
nonlinear feature of G′ as a function of shear strain for
colloidal gels which yields
∞
G′(γ) ∝
1
I4n+5I2n+2
∑
n)0(2n + 1)!
( )
Aγ
2
2n
(4)
k
where Ik ) ∫2π
0 sin θ dθ and A ) (1 + db)/(db - 1), db
being the connectivity or chemical dimension, which
characterizes the scaling of the contour length within
the cluster. This relation has been tested in Figure 2,
inset (a), where all reduced storage moduli G′/G′0
obtained at different temperatures have been plotted
as a function of shear strain. A second-order power
Figure 4. Storage modulus (full symbol) and loss modulus
(open symbol) of the 4 wt % solution as a function of frequency
at 12.5 (9, 0), 25 (b, O), and 50 °C (2, 4).
series of eq 4, i.e., n ) 0, 1, 2 with db ≈ 2.5, matches
remarkably well the shear strain dependence of the G′
modulus up to γ ∼ 1.5, i.e., well above the linear
viscoelastic range. This result suggests that the connectivity of the network is self-similar in this lowest
range of strain amplitude for all tested temperatures.
Moreover, it has to be noted that the strain hardening
strength, denoted as G′max/G′0 and plotted in inset (b)
as a function of temperature, increases up to T ∼ 30 °C
where it starts to decrease gradually with increasing
temperature. According to the molecular approach of the
strain hardening, the increase of G′max/G′0 upon heating
until T ∼ 30 °C suggests (i) an increase of intermolecular
interactions, which leads to an increase of the number
of the elastically active chains, and/or (ii) the existence
of a more stretched conformation of the polymer backbone. Above T ∼ 30 °C, the shear strain amplitude
denoted as γmax, associated with the maximum in G′,
decreases sharply upon heating, reflecting a rupture of
the transient network at weaker deformations when the
temperature is increased. A direct consequence of this
effect is the progressive decrease of the strength of the
strain hardening of G′ above T ∼ 30 °C. The selfsimilarity of the network structure observed at low
strain amplitude could reflect the predominance of the
stretching effect while intermolecular associations may
be favored at higher strain amplitudes, leading to a
more structured network.
As far as the loss modulus G′′ is concerned, its strain
hardening strength increases continuously with increasing temperature. On a molecular level, the loss modulus
reflects generally the effective volume occupied by the
transient network. Consequently, the strain hardening
of the loss modulus may originate from the extension
(stretching) of the polymer coil that results in an
increase of the volume occupied by the network. The
above arguments might suggest that the length of the
polymer coil increases continuously with increasing
temperature.
Figure 4 shows the linear viscoelastic behavior of the
solution as a function of frequency at T ) 12.5, 25, and
50 °C. At T ) 12.5 °C, G′ and G′′ moduli are respectively
proportional to ω2 and ω1 at low frequencies, corresponding to the classical terminal zone. At T ) 25 °C,
the G′ modulus is significantly higher than the G′′
modulus, and both moduli increase in parallel slowly
at high frequencies. Such a behavior is characteristic
of fractal gel structures, in agreement with the discussion following eq 4. At 50 °C, the linear viscoelastic
behavior of the solution is qualitatively similar to that
2886
Bossard et al.
Figure 5. Characteristic time τc, corresponding to the inverse
of the frequency associated with the G′-G′′ crossover, as a
function of temperature.
Figure 6. Apparent viscosity of the 4 wt % solution as a
function of shear stress at T ) 12.5 (2), 18 (9), 22 (b), 35 (1),
and 50 °C (tilted 2).
obtained at 25 °C, but both moduli are globally enhanced by increasing temperature. Moreover, the G′G′′ intersection is observed at a higher characteristic
frequency, ωc, than at 25 °C.
In Figure 5 the characteristic time τc ) 2π/ωc determined at the G′-G′′ intersection is plotted as a function
of temperature. For T < 20 °C, the order of magnitude
of the characteristic time of about 2 s is in good
agreement with that classically encountered for associative polymers for which association/dissociation processes dominate.18,19 A sudden increase of the characteristic time appears clearly at T ∼ 20 °C. Such a jump,
similar to a characteristic time divergence, resembles
a sol/gel transition. Above 30 °C, the characteristic time
decreases continuously with increasing temperature.
This result shows that the molecular dynamics become
faster by increasing temperature above T ∼ 30 °C, which
could be partly attributed to the enhancement of the
thermal motion of polymeric molecules. However, the
latter cannot justify solely the magnitude of the observed τc decrease, and probably other changes of the
interactions sensitive to temperature may occur.
b. Steady Shear Measurements. Figure 6 shows
the apparent viscosity of the polymer solution as a
function of shear stress for temperatures between 12
and 50 °C. Three different temperature regimes can be
identified from this figure. Below T ∼ 20 °C, the flow
curves exhibit a Newtonian plateau η0 at low shear
stress followed by a shear thinning above a critical shear
stress. In this temperature regime, an increase in
temperature induces an increase of the Newtonian
viscosity associated with a sharp decrease of the critical
shear stress value. This behavior reveals a gradual
Figure 7. Newtonian viscosity η0 and intensity of the shear
thickening effect (ηmax - η0)/η0 expressed in percentage of the
Newtonian viscosity of the 4 wt % solution as a function of
temperature.
structure formation in the solution upon heating. The
viscous behavior changes above the temperature T ∼
20 °C that corresponds to the abrupt τc enhancement
evidenced from viscoelastic measurements. In particular, a shear thickening effect appears just after the
linear regime, followed by an abrupt shear thinning.
Such discontinuity in the flow curve is generally observed in the viscous response of a physical gel, for
which the intermolecular dissociation process prevails
beyond the critical shear stress, in the dissociation/
association competition. However, this high-temperature regime (T > 20 °C) can be divided in two subregimes. Between T ∼ 20 and 35 °C, the Newtonian
viscosity and the critical shear stress still increase with
temperature while they both decrease gradually above
T ∼ 35 °C. The temperature dependence of the Newtonian viscosity η0 and the strength of the shear thickening effect quantified by (ηmax - η0)/η0, where ηmax is the
maximum value of the viscosity, are depicted in Figure
7. The Newtonian viscosity profile points out three
temperature regimes similar to those observed previously. The temperature at T ∼ 20 °C marks both a
significant reduction of the rate of increase of the
Newtonian viscosity upon heating and a sudden increase of the shear-thickening effect. On the contrary,
above T ∼ 35 °C the Newtonian viscosity decreases
according to an Arrhenius law
η0 ∼ exp(Ea/RT)
(5)
where the activation energy Ea ∼ 46 kJ/mol can be
considered as the potential barrier to disengage a chain
from a junction point. The value of Ea is in the same
order of magnitude but slightly lower than that determined for hydrophobically associated HEUR telechelic
polymer with similar weight-average molecular weight
end-capped by C16H33O hydrophobic groups.20,21 The
thermothinning behavior observed above 35 °C, a classical behavior for complex and/or simple fluids, like most
polymer solutions, is due to the gradual increase of the
macromolecular thermal motion with increasing temperature which has been previously observed in Figure
5. The increase of the thermal motion might also be
probably responsible for the progressive vanishing of the
shear-thickening effect.
Dynamic Light Scattering. Figure 8 illustrates
representative experimental data, i.e., correlation functions C(q,t) ) x(g(2)(q,t)-1)/f* at a scattering angle θ )
90° (q ) 0.022 nm-1) for a dilute, 0.5 wt %, polymer
Triblock Polyampholyte in Aqueous Solutions 2887
Figure 9. Temperature dependence of the hydrodynamic
radii, RH, of the slowest relaxation modes exhibited in the
intensity autocorrelation function of the 0.5 wt % polymer
solution.
Figure 8. Representative autocorrelation functions, C(q,t), of
the 0.5 wt % polymer solution at θ ) 90 °C (q ) 0.022 nm-1)
and three temperatures, T ) 10, 25, and 40 °C. Open symbols
stand for the experimental data. Solid lines through the data
points represent the best fit results using eq 2. Dashed lines
correspond to the distribution of relaxation time, L(ln τ),
obtained from inverse Laplace transformation. Inset: q dependence of the diffusion coefficient D ) Γ/q2 ) 1/q2τ for the
three relaxation modes.
solution (c/c* ) 0.2) at three temperatures, as well as
the corresponding distributions of relaxation times
obtained with the aid of the inverse Laplace transform
(ILT) technique (cf. eq 2). The solid line through the
experimental points (open symbols) is the best fit curve
using eq 2.
At first sight, the correlation functions seem to exhibit
a two-step relaxation pattern with well-separated fast
and slow modes. Analyzing the experimental data with
a double stretched exponential formalism, we found that
the fast mode is purely exponential while the slow
process is stretched with stretching exponent of about
0.5. This appreciable stretching implies either polydispersity or the existence of species with different sizes.
Indeed, the ILT distributions revealed the existence of
two kinds of particles whose hydrodynamic radii, as
calculated by means of eq 3, correspond to aggregates.
All these three modes are found to exhibit diffusive
character as evidenced from the q independence of the
diffusion coefficient D ) Γ/q2 shown in the inset of
Figure 8. The decay rates Γ were calculated from the
maxima of the ILT distributions. The fast mode corresponding to a “particle” size of about 1 nm should be
associated with ion diffusion observed in pure P2VP at
the same conditions.22 The other two modes correspond
to sizes of about 26 and 170 nm, respectively. The
temperature dependence of the hydrodynamic radii of
the two slow modes is shown in Figure 9. Since the
radius of gyration, Rg, of single chains in salt-free
solution is estimated to be 11 nm,22 the two slow modes
reflect the presence of small associates (with 2-3
associated chains) and large clusters, respectively.
Direct observation by atomic force microscopy confirms
that in this concentration regime the majority of single
chains are participating in small assemblies and larger
loose clusters.23
The temperature dependence of the hydrodynamic
radii exhibits an interesting behavior. With increasing
temperature in the range 10-25 °C the hydrodynamic
radii of the two slow modes, small associates and
clusters, grow from 23.3 to 30 nm for the former and
from 166 to 178 nm for the latter. This effect is
consistent with the viscosity data which show also a
drastic increase in the same temperature range. Increasing further the temperature to 40 °C, we observe
a modest speed-up of the diffusion associated with these
two modes or equivalently a decrease in the related
hydrodynamic radii; this fact again reflects viscosity
changes above 25 °C. It is obvious that both of the
apparent hydrodynamic radii increase with temperature
and reach a maximum value at T ) 25 °C. At low
temperatures (T < 15 °C), the PAA outer blocks adopt
a compact coil conformation since they are close to theta
conditions (UCST).24 On the other hand, the main P2VP
central block is partially protonated and therefore
exhibits a more stretched conformation.25 In such a
situation, intermolecular association through electrostatic interactions and/or hydrogen bonding are prevented, and the viscosity of the system is low. The
swelling of the polymer chain observed upon heating
should be mainly attributed to the expansion of the PAA
coils (excluded-volume effect) at both ends of the macromolecule giving rise to the development of the intermolecular interactions and thus to an increase of the
elastically active chains which contributes to the storage
modulus enhancement. It is worth mentioning that
molecular dynamics simulations have predicted a reversible swelling of neutral random polyampholyte
backbone characterized by a hysteresis when the longrange Coulomb force and short-range attraction force
cooperate.26,27 Such specific predisposition for polyampholytes could explain the hysteresis observed in dynamic moduli of solutions subjected to a temperature
cycle.
Dynamic light scattering results support the analysis
obtained from the rheological results, which suggests a
more effective intermolecular association of the polymer
upon heating. The molecular swelling of the PAA outer
blocks favors intermolecular interactions mainly between the oppositely charged blocks,14 leading to the
formation of a physical gel. However, most of the PAA
units and about 70% of P2VP units (pH 3.4) are
uncharged, and hydrogen bonding between the different
moieties should exist which stabilize the association.28
Simultaneously, the thermal motion increases and its
2888
Bossard et al.
effect becomes predominant above T ∼ 35 °C, for which
the molecular expansion reaches a certain limit. The
thermal motion prevalence at high temperature is
related with the significant decrease of the characteristic time determined by dynamic rheology and the
weakening of the shear-thickening and the strainhardening effects. However, the magnitude of this
decrease cannot be ascribed only to thermal motion,
suggesting a weak alteration of the structure which is
evident by light scattering. The observed decrease of the
size of the clusters could be attributed to the weakening
of the hydrogen-bonding contribution on the intermolecular association since the H-bonds are not favored
upon heating.29
Concluding Remarks
In this study, a rich and rather unexpected thermosensitivity of an asymmetric triblock polyampholyte
of the type PAA135-P2VP628-PAA135 in salt-free aqueous solutions has been presented through rheological
measurements and supported by dynamic light scattering. With increasing temperature, the whole set of
rheological data demonstrate a sol/gel like transition,
2 orders of magnitude viscosity enhancement, while the
rheological properties of the gel above T ∼ 35 °C exhibit
Arrhenius behavior. This peculiar thermal response
results from the competition between the significant
swelling of the PAA outer blocks, which favors intermolecular interactions responsible for the pronounced
thermothickening behavior and the thermal motion,
which weaken the rheological properties of the polymer
solution by speeding up the molecular dynamics. The
partial expansion of the polymer chains upon heating
is a consequence of the enhancement of the solvent
quality. The results reported in this study emphasize
the serious impact of the solvent quality in the rheological behavior of soluble polymers, which is frequently
overlooked or taken into account less seriously.
The novelty of this thermoresponsive behavior arises
from the fact that none of the polymeric components of
the copolymer exhibit LCST, which was the only reason
known so far to induce the thermothickening effect in
associative polymers. On the contrary, the outer PAA
blocks exhibit UCST. In such a case, coil expansion
occurs upon heating, allowing the intermolecular interactions mainly among oppositely charged moieties to
develop, leading eventually to the formation of an
infinite transient network.
Acknowledgment. We thank Prof. Dimitris Vlassopoulos, Prof. Thierry Aubry, and Prof. G. Staikos for
fruitful discussions and comments about this work,
which has been performed with the financial support
of the European Community under Grant HPMDCT2000-00054-02.
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Bossard, F.; Sfika, V.; Tsitsilianis, C. Macromolecules 2004,
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(15) Berne, B. J.; Pecora, R. Dynamic Light Scattering with
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Correlation Techniques in Fluid Mechanics; Schulz-DuBoir,
E. O., Ed.; Springer-Verlag: Berlin, 1983; p 15.
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Macromolecules 1998, 31, 4149.
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MA047520P
PAPER
www.rsc.org/softmatter | Soft Matter
pH-Tunable rheological properties of a telechelic cationic polyelectrolyte
reversible hydrogel
Frédéric Bossard,a Thierry Aubry,a Georgios Gotzamanisb and Constantinos Tsitsilianis*b
Received 31st January 2006, Accepted 18th April 2006
First published as an Advance Article on the web 3rd May 2006
DOI: 10.1039/b601435f
Steady shear properties and linear and nonlinear viscoelastic behaviors of a poly(methyl
methacrylate)–poly(dimethyl amino ethyl methacrylate)–poly(methyl methacrylate) polymer,
(PMMA–PDMAEMA–PMMA), telechelic polymers in salt-free aqueous solution have been
investigated as a function of concentration and pH. Above a critical concentration, a transient
physical network is formed through an association mechanism between hydrophobic end groups,
leading to a gel-like behavior. The gel-like polymer solutions were shown to exhibit a peculiar flow
behavior, associated with time fluctuation of the transient first normal stress difference, attributed
to orientation effects of the stiff charged polymer chains. The viscoelastic behavior was shown to
be governed by two pH dependent time scales: a short time scale controlled by the lifetime of the
hydrophobic associative junctions and a long time scale corresponding to the network relaxation
time. All rheological results show strong evidence that Coulomb interactions, which control both
macromolecular chain rigidity and inter-chain interactions, lead to specific pH-tunable properties
of great potential interest.
I. Introduction
Traditional hydrophobically modified water-soluble polymers
are hydrophilic neutral macromolecules bearing short hydrophobic groups that self-associate in water, leading to the
formation of a physical transient network above a threshold
concentration.1 The location of the hydrophobic groups is
known to play an important role on the rheological and
structural properties of associative polymers. They can be
either distributed along the hydrophilic backbone or located at
both ends; in the latter case they are named telechelic
polymers. Rheological and structural properties of traditional
neutral telechelic associative polymers have been explored in
experimental,2,3 theoretical4,5 and computational6,7 studies.
At a critical micelle concentration, hydrophobic end groups
self-associate in a compact core surrounded by hydrophilic
central polymeric chains adopting loop conformations in
flower-like micelles.8,9 On increasing polymer concentration,
some hydrophobic chain ends can disengage from micelles
leading to the formation of bridges between neighboring
micelles, and eventually to the formation of a threedimensional physical transient network in which flower-like
micelles act as nodes. The flow behavior of this network is
characterized by a Newtonian viscosity at low shear rates,
shear-thickening behavior at intermediate shear rates, followed
by shear-thinning behavior at higher shear rates. The
molecular interpretation of these non-linear regimes is still a
matter of debate. For example, the shear-thickening behavior
a
Laboratoire de Rhéologie, 6 avenue Le Gorgeu - CS93837, 29238 Brest
Cedex 3, France
b
Department of Chemical Engineering, University of Patras 26504,
Patras Greece and Institute of Chemical Engineering and High
Temperature Chemical Processes, FORTH/ICE-HT.
E-mail: [email protected]; Fax: +30 2610 997 266
510 | Soft Matter, 2006, 2, 510–516
has been attributed either to a shear-induced increase of
the density of elastically active chains10,11 or to the non-linear
stretching behavior of the connecting polymeric chains, or
to incomplete relaxation of dangling chains.5,12 As far as the
linear viscoelastic behavior is concerned, the response is welldescribed by a single time Maxwell model. The Maxwellian
behavior shows that the dynamics of telechelic polymers is
governed by a single relaxation time, which is related to the
average lifetime of an associative junction.2
In the past few years, a new class of telechelic polymers has
been designed in which the central backbone has a polyelectrolyte character.13–17 The main feature that makes
telechelic polyelectrolytes distinct from conventional telechelics (non-ionic hydrophylic part) is that the chain conformation of the bridging chains of the resulting transient network
depends on its degree of neutralization and can be tuned by
external stimuli such as pH and ionic strength. At a certain pH
(depending on the polyelectrolyte nature) and at free salt
conditions the central chain adopts a stretched conformation,
which has fundamental consequences on the association
mechanism and the rheological properties. For example, the
sol–gel transition appears at a lower concentration, compared
to that for neutral telechelic polymers and the flow curve
exhibits an unusual complex viscosity profile characterized by
a yield stress and several shear-thinning regimes.13–15 Recently,
the yield stress was shown to be apparent, as a Newtonian
plateau, was finally determined at very low shear stresses using
creep measurements.16
Very recently a novel type of telechelic polyelectrolyte composed of a long PDMAEMA end-capped by short PMMA
blocks was designed and its self assembly capability was
explored in aqueous media.17 The MMA32–DMAEMA224–
MMA32 copolymer self-associates into micelles through
hydrophobic interactions of the PMMA end-blocks. On a
This journal is ß The Royal Society of Chemistry 2006
second level of hierarchy, a three dimensional transient
network (Fig. 1) forms, leading eventually to a stiff physical
gel in salt-free aqueous solutions, above 1 wt% polymer
concentration. The novel character of the self-assemblies is
that the main PDMAEMA hydrophilic part of the copolymer
adopts a stretched conformation due to protonation of its
monomer-repeating units. Therefore, as can be seen in Fig. 1,
the repulsive electrostatic interactions along the hydrophilic
chains together with their relative short length, i.e. small
number of Kuhn segments, prevent the formation of loops,
which gives specific features to the topology of the network
and of the micelles, which adopt a star-like shape.17
In this work, we present the rheological properties of the
MMA32–DMAEMA224–MMA32 cationic telechelic polyelectrolyte, in salt-free aqueous media. The concentration
and pH dependence of the linear and non-linear rheological
properties was thoroughly studied. Particular attention was
paid to pH dependence in order to offer a physical insight
into the various effects of Coulomb interactions on the
rheological properties of telechelic polyelectrolyte associative
polymers. The results were discussed in terms of microstructures and compared mainly with those of traditional
neutral telechelic associative polymers.
The remarkable specific rheological properties of the
cationic telechelic polyelectrolyte physical gels found herein,
combined with the relative good biocompatibility (cell viability
50–60%) and capability of the protonated PDMAEMA
cationic major component of the polymer to be complexed
with DNA, could make this polymer a good candidate for
research towards biological applications.18
II. Experimental
than 4 and decreases sharply as pH increases, being negligible
at pH > 9. The so-synthesized MMA32–DMAEMA224–
MMA32 copolymer is water-soluble in a wide range of pH
but precipitates at pH higher than 10.86. Polymer solutions
were obtained by dissolving the proper amount of polymer in
distilled water, and pH adjustment was achieved by adding a
negligible amount of 1 N hydrochloric acid. Rheological
characterizations have been carried out 24 h after the sample
preparation.
(b) Rheometry
The linear and nonlinear rheological measurements were
carried out using two rotational rheometers: a strain controlled
Rheometric Scientific ARES rheometer, equipped either with a
cone and plate geometry (diameter = 50 mm, cone angle =
2.3u, truncation = 48 mm) and striated parallel-plate geometry
(diameter 20 mm, gap 2 mm) to eliminate any wall slip
effects for highly viscous solutions, and a stress controlled
Rheometric Scientific SR 200 rheometer, equipped with a cone
and plate geometry (diameter = 25 mm, cone angle = 5.7u,
truncation = 56 mm). The apparent viscosity versus time was
measured at each stress and the steady-state viscosity value
was determined as the limit, on long time scales, of the
transient viscosity, following the criterion: time evolution of
the transient viscosity lower than 1% during 1 min. A thin
layer of low viscosity silicone oil was put on the air–sample
interface in order to minimize solvent evaporation. After each
sample loading, normal stress was monitored and was shown
to relax within 5 min. Consequently, a delay of 5 min was
applied prior to any measurement, in order to erase the
mechanical history. The temperature, fixed at 25 ¡ 0.1 uC was
controlled by a water bath circulator.
(a) Material
The polymer used in this study is a triblock copolymer
composed of a long central poly(dimethyl amino ethyl
methacrylate) (PDMAEMA) chain end-capped by short
poly(methyl methacrylate) (PMMA) blocks. This telechelic
copolymer was synthesized by a ‘‘living’’ polymerization
method called group transfer polymerization, GTP. Details
of the synthesis were previously reported.17 The degree of
polymerization is 224 for the central PDMAEMA block and
32 for the hydrophobic PMMA blocks, corresponding to a
molecular weight Mn = 4.2 6 104 g mol21 and a molecular
polydispersity Mw/Mn = 1.26, as determined by GPC analysis.
The ionization degree of the central DMAEMA224, measured
by hydrogen ion titration, reaches more than 90% at pH lower
Fig. 1 Schematic representation of the association mechanism of
MMA32–DMAEMA224–MMA32 cationic telechelic polyelectrolyte in
salt-free aqueous media as deduced from direct AFM observation
(ref. 16). The concentration increases from the left to the right.
III. Results and discussion
The linear and non-linear rheological behavior of the
MMA32–DMAEMA224–MMA32 cationic telechelic polyelectrolyte solutions has been studied as a function of polymer
concentration and pH.
(a) Rheological properties as a function of polymer concentration
The dependence of rheological properties on polymer concentration, c, was investigated at pH 3, at which central
PDMAEMA blocks are significantly protonated. Fig. 2 shows
the apparent viscosity as a function of shear stress for 0.06,
0.3, 0.7, 1, 1.4 and 1.8 wt% polymer solutions. For all polymer
solutions, we ensured systematically that no wall slip effects
occurred, by comparing rheological data obtained using a
striated parallel-plate geometry with various gaps. All polymer
solutions exhibit a linear behavior at low shear stresses,
characterized by a Newtonian viscosity g0. The concentration
dependence of both Newtonian viscosity and specific viscosity
gsp, plotted in Fig. 3, clearly shows three concentration
regimes, characterized by power law functions with distinct
exponents. For c , 0.1 wt%, the power law exponent is
found to be equal to 0.5 for g0 , respectively 2.5 for gsp; in
this concentration regime, macromolecules have already selfassembled into micelles.17 In the intermediate concentration
Soft Matter, 2006, 2, 510–516 | 511
Fig. 2 Apparent viscosity versus shear stress for 0.06, 0.3, 0.7, 1, 1.4
and 1.8 wt% polymer solutions at pH 3.
Fig. 3 Newtonian viscosity ($) and specific viscosity (&) of polymer
solutions at pH 3 as a function of concentration.
regime, 0.1 wt% , c , 1 wt%, g0 and gsp increase sharply, with
a power law exponent equal to 8 for both g0 and gsp; such a
high exponent suggests that, in this concentration regime, a
three dimensional transient network is built through extensive
association of the PMMA dangling hydrophobes.19 In the
upper concentrated regime, where the network is fully
developed, the viscosity turns concentration independent.
The description of three concentration regimes is also
relevant regarding the flow behavior. The lower concentration
regime is characterized by a Newtonian behavior over the
whole range of shear stress investigated, whereas strong
nonlinear behaviors appear for polymer concentrations
c > 0.1 wt%. In the intermediate concentration regime, the
nonlinear behavior is characterized by a shear-thinning
behavior, which is more marked as polymer concentration
increases.
In the upper concentration regime, polymer solutions
exhibit a peculiar complex nonlinear behavior. Fig. 4 shows
the apparent viscosity and the first normal stress difference N1
as a function of shear stress for a 1 wt% polymer solution. All
viscosity values have been determined as steady state limits of
the transient response. The flow curve clearly exhibits four
512 | Soft Matter, 2006, 2, 510–516
Fig. 4 Apparent viscosity and first normal stress difference N1 of a
1 wt% polymer solution at pH 3.5 as a function of shear stress. Inset:
time dependence of the first normal stress difference at shear stress of
20 Pa and 60 Pa.
distinct regions: a zero-shear Newtonian region (0), a
Newtonian region (II) at intermediate shear stresses, which
separates two shear-thinning regions (I) and (III) at low and
high shear stresses respectively. The shear-thinning region (I)
corresponds to a sharp four decades decrease of viscosity level,
usually attributed to an apparent yield stress, whereas a
smoother shear-thinning behavior is observed in the high shear
stress Region (III). The transient first normal stress difference
presented in inset of Fig. 4 exhibits time fluctuations in regions
(0), (I) and (II), where no steady-state value can ever be
measured, whereas a steady state regime is finally achieved
after a chaotic transient regime in region (III). To our
knowledge, time fluctuations of N1 have never been observed
for associative polymeric systems; we suggest attributing these
oscillations to the stiffness of the macromolecular backbone,
due to the electrostatic intra-molecular repulsive interactions.
Indeed this peculiar time dependence of N1 is reminiscent of a
time-periodic fluctuation of the direction of average molecular
orientation of stiff molecules between two limiting angles
in a shearing flow.20 Region (II) then appears as an
intermediate region between region (I) characterized by the
time-periodic fluctuation of the direction of average molecular
orientation, and region (III), corresponding to a progressive
shear-induced alignment of the direction of the average
molecular orientation. In Region III, the first normal stress
difference increases slightly with increasing shear rate according to a power law with an exponent 0.5, that is much lower
than the exponent 2, generally obtained for polymer solutions
at low shear rates.
In order to complete the rheological characterization,
the viscoelastic behavior of a 1 wt% polymer solution was
investigated. Fig. 5 shows the storage modulus G9 and the loss
modulus G0, at a pulsation v = 1 rad s21, as a function of the
amplitude of shear strain c0, at pH 3.5. The polymer solution
exhibits a linear viscoelastic response below a critical shear
strain cc y 0.7%, where G9 > G0. Above this narrow linear
viscoelastic regime, G9 decreases as shear strain increases
whereas G0 increases, reaches a maximum for c0 y 3%, and
decay and a stretched exponential function that depicts long
time relaxation.
n t
t
Gðt,cÞ~G0,f ðcÞ exp {
zG0,s ðcÞ exp {
(1)
tf (c)
ts ðcÞ
Fig. 5 Storage modulus and loss modulus of a 1 wt% polymer
solution at pH 3.5 as a function of shear strain.
Fig. 6 Stress relaxation modulus of a 1 wt% polymer solution at
pH 3.5 for strain amplitudes c ranging from 0.4% to 20%. The
continuous lines through data represent the best-fitting calculation
curves obtained from eqn (1).
then decreases. Such a G0 feature at intermediate strain
amplitudes was observed with other telechelic polyelectrolytes15 and was ascribed to the strain induced imbalance
between junction destruction rate and junction creation rate
within the network.21
The 1 wt% polymer solution at pH 3.5 was submitted to
stress relaxation measurements in the linear (c = 0.4%) and
non-linear (c = 1.5%, 2%, 2.5%, 10% and 20%) viscoelastic
regime. Stress relaxation functions are plotted in Fig. 6. The
relaxation functions G(t) exhibit a two-step relaxation pattern,
proving the existence of a fast and a slow relaxation mode.
Experimental data are properly fitted by eqn (1), a sum of a
mono-exponential function that describes the fast initial
where G0 is the instantaneous modulus and t a characteristic
relaxation time; the indices f and s referring to ‘‘fast’’ and
‘‘slow’’ relaxation modes respectively. The exponent 0 , n , 1
is the stretched exponent that quantifies the departure from the
mono-exponential function; it measures the broadness of the
time relaxation distribution, the smaller n values corresponding to the broader distributions.
It has to be noticed first that the weak short time
exponential relaxation at low strains makes the fast mode
quite difficult to observe in the linear viscoelastic regime. The
best fit is obtained for a value of tf y 0.5 s in the linear
viscoelastic regime, which decreases as strain increases, reaching 0.12 s at c = 20%. On the contrary, the long relaxation
time ts is about 90 s in the linear viscoelastic regime and
increases with strain, reaching 186 s at c = 20%. Additionally,
the stretched exponent has a value of y 0.7 in the linear
viscoelastic regime, and decreases with increasing strain,
reaching y 0.4 at c = 20%, showing a significant broadening
of the relaxation time distribution with increasing deformations. Table 1 presents the rheological parameters corresponding to the best fit of experimental results using eqn (1).
A two-step relaxation mechanism was observed and
discussed by Séréro et al. for neutral telechelic associating
polymers in the nonlinear regime.22 They attributed the twostep relaxation to the existence of two populations of
elastically active chains: highly stretched chains and weakly
stretched chains. There are still two major differences in
relaxation response between the system studied by Séréro et al.
and that studied in the present paper, which can be attributed
to the different nature (ionic/neutral) of the bridging hydrophilic chains. The first difference is that the two relaxation
processes appear only at very large strains c > 200%, well
above the linear viscoelastic range in the case of neutral
telechelic polymers, whereas they appear at very low strains,
that is just above cc y 0.7%, in the case of charged telechelic
polymers. The second significant difference appears in the
strain dependence of the two characteristic relaxation times.
Indeed both relaxation times decrease with increasing strain in
the case of neutral telechelic polymers, whereas the relaxation
times have opposite strain dependence in the case of charged
telechelic polymers: the short relaxation time decreases and the
long relaxation time increases with increasing strains.
From a molecular point of view, the fast relaxation
mechanism may be attributed to polymer chains that relax
Table 1 Rheological parameters G0,f, G0,s, G0,s/G0,f, tf, ts and n obtained from the fit of stress relaxation modulus of a 1 wt% polymer solution at
pH 3.5 using eqn (1)
Strain (%)
G0,f/Pa
0.4
2
2.5
10
20
150.1
105.9
114.1
102.2
69.1
¡
¡
¡
¡
¡
G0,s/Pa
1.5
1.2
1.6
1.9
1.1
120.3
75.5
61.7
17.8
10.0
¡
¡
¡
¡
¡
0.4
0.4
0.2
0.2
0.1
G0,s/G0,f
tf/s
ts/s
0.80
0.71
0.54
0.17
0.14
0.52 ¡ 0.08
0.41 ¡ 0.05
0.30 ¡ 0.03
0.1 ¡ 0.06
0.12 ¡ 0.07
90.4
102.2
126.4
151.6
186.1
n
¡
¡
¡
¡
¡
1.1
2.1
1.3
2.5
2.1
0.7 ¡ 0.1
0.6 ¡ 0.06
0.63 ¡ 0.01
0.51 ¡ 0.10
0.4 ¡ 0.05
Soft Matter, 2006, 2, 510–516 | 513
following a process controlled by the disengagement of the
hydrophobic blocks from associative junctions. Thus this
mechanism is essentially comparable to that governing the fast
relaxation in neutral telechelic polymers, as suggested by the
similarity of the strain dependence of the short relaxation time
for both systems. However some differences between charged
and neutral telechelics in the molecular dynamics of the fast
relaxation process are expected. Indeed intra-chain electrostatic repulsive interactions increase chain stiffness, which
may induce two different coupled opposite effects: a decrease
of the short relaxation time with increasingly stretched
conformation, coupled with an increase of the short
relaxation time due to an increase of the effective association
lifetime. The latter effect proposed by Rubinstein and
Semenov,19 is due to the difficulty for a given hydrophobic
end group, which disengage from an associative junction, to
find a new available partner; this effect could also be the result
of the inter-chain repulsions.
The slow relaxation mechanism may be attributed to
polymer chains, which are elastically active in the network
and cannot relax on short times because of electrostatic
interactions with neighboring chains. We suggest that these
interactions are either electrostatic repulsions between positively charged central blocks, or dipole attractions arising
from counterion condensation on the polyelectrolyte chains.23
Such electrostatic interactions act as ‘‘electrostatic entanglements’’, which tends to slow down the chain relaxation. The
increase of the long relaxation time with increasing strain is
attributed to the strain-induced reinforcement of the number
and/or the intensity of these interactions.
Thus the presence of two distinct relaxation time scales is
attributed to a network structure, with a non homogeneous
complex topology, composed of two polymer chain populations: chains whose relaxation is governed by the effective
lifetime of an associative junction (fast relaxation), and chains
whose relaxation is hindered by electrostatic inter-chain
interactions (slow relaxation). The number density of the two
populations within the network is a function of strain, as
shown by the strain dependence of the ratio of the two
instantaneous relaxation moduli G0,s over G0,f, in Table 1.
More precisely, the decrease of the slow relaxation contribution to the total instantaneous relaxation modulus with
increasing strain suggests that the strain induces an increase
of the number density of the elastically inactive polymer chains
that relax according to the fast relaxation mechanism.
(b) Influence of pH
In this section, pH-dependent properties of the cationic
telechelic polyelectrolytes have been investigated performing
steady shear and dynamic oscillatory measurements.
In Fig. 7 (a) the zero-shear viscosity of a 1 wt% polymer
solution has been plotted as a function of pH. The viscosity
passes through a maximum, of about 50000 Pa, in the vicinity
of pH 4. Close to pH 4, the degree of ionization of the
PDMAEMA chains is about 90 %, which results in stretched
macromolecular conformation.17 At pH > 4, increasing pH
leads to a gradual deprotonation of the monomeric units,
which lowers the degree of ionization and increases chain
514 | Soft Matter, 2006, 2, 510–516
Fig. 7 (a) pH dependence of zero shear viscosity of a 1 wt% polymer
solution. A stiff gel is formed at pH 4 (middle digital photograph).
(b) Viscosity profile of a 1 wt% polymer solution at pH 1(&), 4($)
and 7(m).
flexibility, favoring bridge to loop transitions and hence
viscosity decrease. On the other hand, at pH , 4, the ionic
strength of the solution increases upon lowering pH, leading to
electrostatic screening effects that result in viscosity decrease.
Moreover, the apparent yield stress, that is the stress at which
viscosity departs from the low shear Newtonian plateau,
has also a maximum value, about 30 Pa, in the vicinity of
pH 4, as shown in Fig. 7 (b). At last, it has to be noticed that
regions I and II of the flow curves, plotted in Fig. 7 (b), are
more marked at pH 4. All these linear and non-linear
rheological properties, studied as a function of pH, suggest a
soft–stiff–soft gel smooth transition around pH 4.
The storage and loss moduli of the 1 wt% polymer solution
at pH 3.5, 5.5 and 7.5 are plotted in Fig. 8, in order to study
the effect of pH on the two previously defined characteristic
time scales governing the linear viscoelastic response. First of
all, Fig. 8 shows that all polymer solutions exhibit a similar
linear viscoelastic behavior in the range of pH explored.
However, the Cole–Cole plots, inserted in Fig. 8, show that
the two relaxation time scales, previously observed in stress
relaxation experiments, are clearly pH dependent. The slow
relaxation mode, corresponding to the first maximum of the
Fig. 8 Storage modulus (full symbols) and loss modulus (open
symbols) of a 1 wt% polymer solution at pH 3.5 (&, %), pH 5.5
($, #) and pH 7.5 (m, n). Inset: Cole–Cole plot at pH 3.5 (%), pH 5.5
(#) and pH 7.5 (n). The dash line represents the Cole–Cole curve of a
single Maxwell element.
curve, is characterized by a wide spectrum of relaxation times,
as ascertained by the significant departure of the Cole–Cole
plots from the semicircle. The fast relaxation mode is not so
well defined in the Cole–Cole plots; it is simply suggested by
the significant upturn in G9 and G0 moduli at high frequencies.
From Fig. 8, it is clear that decreasing pH from 7.5 to 5.5
results in increasing both relaxation times. This effect can be
attributed to an increase of the charge density (degree of
ionization) of central PDMAEMA blocks as pH decreases.
Indeed, increasing charge is expected to have two effects:
- it increases chain stiffness and therefore the effective
association lifetime controlling the fast relaxation mode, as
discussed previously.
- it favors counterion condensation, which enhances the
electrostatic inter-chain attractions between central blocks of
elastically active chains, therefore increasing the slow relaxation time.
IV. Conclusion
In this paper, the influence of polymer concentration and
pH on the rheological behavior of a cationic PMMA–
PDMAEMA–PMMA telechelic polymer was thoroughly
investigated. Above a critical concentration, which depends
on pH, polymer chain-ends interact via hydrophobic interactions and form a transient physical network exhibiting peculiar
linear and non-linear rheological behaviors. The specificity of
the rheological properties, compared to those of neutral
telechelic polymer solutions, is due to intra- and inter-chain
Coulombic repulsive and/or attractive interactions.
In steady shear flow, the flow curves are quite similar to
those obtained with telechelic polymers, but the non-linear
elastic properties exhibit a quite peculiar feature, a time
fluctuation of the transient first normal stress difference,
which was attributed to orientation effects of the stiff charged
polymer chains.
As far as the viscoelastic behavior is concerned, telechelic
polyelectrolyte solutions exhibit a linear response governed by
two distinct characteristic relaxation time scales, corresponding to two chain populations. The short relaxation time scale is
that of polymer chains whose relaxation is controlled by the
effective disengagement time of a hydrophobic group from an
association. The long relaxation time scale is that of elastically
active polymer chains of the network, interacting with
neighboring chains via repulsive and/or attractive electrostatic
interactions. Both relaxation time scales were pH dependent
and oppositely dependent on strain in contrast with the
behavior of non-ionic associative telechelics. Finally a maximum in some relevant rheological material functions were
observed at pH 4.
To our knowledge, this is the first paper presenting a
thorough experimental investigation of the influence of pH
driven electrostatic interactions on the rheological properties
of positively charged cationic telechelic polymer networks.
Such a novel biocompatible and pH responsive polymeric
system offers many potential applications in pharmaceutic and
cosmetic industries.
At last, the experimental results presented in the paper
could be used to validate any forthcoming model of transient
networks of charged associating systems.
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11252
Langmuir 2007, 23, 11252-11258
Characterization of the Core-Shell Nanoparticles Formed as Soluble
Hydrogen-Bonding Interpolymer Complexes at Low pH
Maria Sotiropoulou,† Frederic Bossard,‡,# Eric Balnois,§ Julian Oberdisse,|,⊥ and
Georgios Staikos*,†
Department of Chemical Engineering, UniVersity of Patras, GR-26504 Patras, Greece, Institute of
Chemical Engineering and High Temperature Chemical Processes, FORTH/ICE-HT, P.O. Box 1414,
26504 Patras, Greece, Laboratoire Polymères, Propriétés aux Interfaces et Composites (L2PIC),
UniVersité de Bretagne Sud, Rue de Saint Maudé, BP 92116, 56321 Lorient, France, Laboratoire des
Colloı̈des, Verres et Nanomatériaux, UMR CNRS/UM2, UniVersité Montpellier II, F-34095 Montpellier,
France, and Laboratoire Léon Brillouin CEA/CNRS, CEA Saclay, 91191 Gif sur YVette, France
ReceiVed May 28, 2007. In Final Form: July 23, 2007
The formation of soluble hydrogen-bonding interpolymer complexes between poly(acrylic acid) (PAA) and poly(acrylic acid-co-2-acrylamido-2-methyl-1-propane sulfonic acid)-graft-poly(N,N-dimethylacrylamide) (P(AA-coAMPSA)-g-PDMAM) at pH ) 2.0 was studied. A viscometric study showed that in semidilute solution a physical
gel is formed due to the interconnection of the anionic P(AA-co-AMPSA) backbone of the graft copolymer, in a
transient network, by means of the complexes formed between the PDMAM side chains of the graft copolymer and
PAA. Dynamic and static light scattering measurements, in conjunction with small-angle neutron scattering measurements,
suggest the formation of core-shell colloidal nanoparticles in dilute solution, comprised by an insoluble PAA/
PDMAM core surrounded by an anionic P(AA-co-AMPSA) corona. Even if larger clusters are formed in semidilute
solution, the size of the insoluble core remains practically stable. Atomic force microscopy performed under ambient
conditions reveal that the particles collapse and flatten upon deposition on a substrate, with dimensions close to the
ones of the dry hydrophobic core.
Introduction
When weak polyacids such as poly(acrylic acid) (PAA) or
poly(methacrylic acid) (PMAA), and proton acceptor polymers,
such as polyethyleneoxide (PEO), or polyacrylamides, are mixed
in solution at pH lower than 3-4, an associative phase separation
takes place,1-7 as a result of the formation of hydrogen-bonding
interpolymer complexes (IPCs). A considerable amount of work
on such hydrogen-bonding IPCs has been presented in two
reviews.8,9 The potential applications of these IPCs to various
fields, such as drug delivery formulations,10-12 biomaterials,13
* To whom correspondence should be addressed. E-mail: staikos@
chemeng.upatras.gr.
† University of Patras.
‡ FORTH/ICE-HT.
§ Université de Bretagne Sud.
| Université Montpellier II.
⊥ CEA Saclay.
# Present address: Laboratoire de Rheologie, UMR 5520, Université Joseph
Fournier, 1301, rue de la piscine, BP 53, 38041 Grenoble Cedex 9, France.
(1) Bailey, F. E.; Lundberg, Jr., R. D.; Callard, R. W. J. Polym. Sci., Part A:
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Langmuir 2004, 20, 3785-3790.
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(12) Carelli, V.; Di, Colo, G.; Nannipieri, E.; Poli, B.; Serafini, M. F. Int. J.
Pharm. 2000, 202, 103-112.
emulsifiers,14 and membrane and separation technology,15,16 has
further stimulated the research interest in this field.
To extend the solubility of the hydrogen-bonding IPCs in the
low pH region, some efforts have recently been undertaken.17,18
In one of them, an anionically charged graft copolymer, poly(acrylic acid-co-2-acrylamido-2-methyl-1-propane sulfonic acid)g-poly(N,N-dimethylacrylamide) (P(AA-co-AMPSA)-g-PDMAM), has been synthesized by grafting poly(N,Ndimethylacrylamide) (PDMAM) chains onto an acrylic acidco-2-acrylamido-2-methy-1-propane sulfonic acid copolymer
(P(AA-co-AMPSA)) backbone. PDMAM is a water-soluble
polymer with important proton acceptor properties, forming
hydrogen-bonding IPCs with PAA,19,20 which precipitate out
from water even at pH values as high as 3.75.17 When these graft
copolymers are mixed with PAA in a low pH (pH < 3.75) aqueous
solution, hydrogen-bonding IPCs between the PDMAM side
chains and PAA are formed. Nevertheless, the presence of the
negatively charged AMPSA units in the graft copolymer backbone
prevents their precipitation.17 Moreover, rheological measurements in a semidilute solution have shown a gel-like behavior
in this low pH region.21 This behavior has been attributed to the
interconnection of the negatively charged backbone chains of
the graft copolymer by means of the hydrogen-bonding inter(13) Chun, M.-K.; Cho, C.-S.; Choi, H.-K. J. Controlled Release 2002, 81,
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(17) Sotiropoulou, M.; Bokias, G.; Staikos, G. Macromolecules 2003, 36,
1349-1354.
(18) Ivopoulos, P.; Sotiropoulou, M.; Bokias, G.; Staikos, G. Langmuir 2006,
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Okano, T. Macromolecules 2000, 33, 444-450.
(21) Bossard, F.; Sotiropoulou, M.; Staikos, G. J. Rheol. 2004, 48, 927-936.
10.1021/la701561y CCC: $37.00 © 2007 American Chemical Society
Core-Shell Nanoparticles as Interpolymer Complexes
polymer complexes formed between the PDMAM side chains
of the graft copolymer and the PAA chains.
Small-angle neutron scattering (SANS) measurements, already
used to show the formation of dense hydrogen-bonding IPCs
between PEO and partially neutralized (3-9%) PMAA,22 were
also used to study the microstructure of the above-mentioned
electrostatically stabilized colloidal system in D2O.23 Coreshell nanoparticles comprised by an insoluble hydrogen-bonding
IPC core and a hydrophilic negatively charged corona surrounding
it was supposed to be formed. The formation of similar colloidal
complexes has also been observed as a result of the interaction
of polyelectrolyte-neutral block copolymers or of comb-type
polyelectrolytes with oppositely charged synthetic or biological
macromolecules24-28 and surfactants.29-33 However, there is a
major difference in our case. The colloidal nanoparticles formed
are pH-sensitive as they are formed at low pH and dissociate at
pH > 3.75.17
In this work we have proceeded to a thorough study of the
interactions between a P(AA-co-AMPSA)-g-PDMAM graft
copolymer, containing 48 wt % of PDMAM, shortly designated
as G48, and PAA, at pH ) 2.0, in a broad concentration region,
ranging in the dilute and the semidilute regime. In such a low
pH insoluble IPCs are formed17 as a result of successive hydrogen
bonds between the carboxylic groups of PAA and the amide
groups of PDMAM.6,34 Nevertheless, the particles formed do
not precipitate but remain in a colloidal form in the solution due
to the anionic backbone of the graft copolymer. Rheology
measurements were used for the determination of a critical
concentration, c*, over which gel formation takes place. Dynamic
light scattering (DLS) and SANS measurements indicated the
formation of core-shell nanoparticles, transformed to bigger
clusters as the concentration increased above c*, with their cores
remaining unchanged. Static light scattering in dilute solution
was used to determine the molecular weight of the isolated
nanoparticles and atomic force microscopy (AFM) to estimate
their size after evaporation of the solvent.
Materials. A sample of PAA (Polysciences), with a nominal
molecular weight of 9.0 × 104 Da, was dissolved in a 0.01 N HCl
solution, dialyzed against water through a cellulose membrane with
a molecular weight cutoff equal to 12 kDa (Sigma), and finally
obtained by freeze-drying.
The monomers, acrylic acid (AA), 2-acrylamido-2-methyl-1propane sulfonic acid (AMPSA) (Polysciences), and N,N-dimethylacrylamide (DMAM) (Aldrich), were used as received. Ammonium
persulfate (APS, Serva), potassium metabisulfite (KBS, Aldrich),
2-aminoethanothiol hydrochloride (AET, Aldrich), and 1-(3-(dim(22) Zeghal, M.; Auvray, L. Europhys. Lett. 1999, 45, 482-487.
(23) Sotiropoulou, M.; Oberdisse, J.; Staikos, G. Macromolecules 2006, 39,
3065-3070.
(24) Bronich, T. K.; Popov, A. M.; Eisenberg, A.; Kabanov, V. A.; Kabanov,
A. V. Langmuir 2000, 16, 481-489.
(25) Harada, A.; Kataoka, K. Science 1999, 283, 65-67.
(26) Maruyama, A.; Katoh, M.; Ishihara, T.; Akaike, T. Bioconjugate Chem.
1997, 8, 3-6.
(27) van der Burgh, S.; de Kaizer, A.; Cohen Stuart, M. A. Langmuir 2004,
20, 1073-1084.
(28) Voets, I. K.; de Kaizer, A.; Cohen, Stuart, M. A.; de Waard, P.
Macromolecules 2006, 39, 5952-5955.
(29) Hervé, P.; Destarac, M.; Berret, J.-F.; Lal, J.; Oberdisse, J.; Grillo, I.
Europhys. Lett. 2002, 58, 912-918.
(30) Berret, J.-F.; Vigolo, B.; Eng, R.; Hervé, P. Macromolecules 2004, 37,
4922-4930.
(31) Nisha, C. K.; Basak, P.; Manorama, S. V.; Maiti, S.; Jayachandran, K.
N. Langmuir 2003, 19, 2947-2955.
(32) Balomenou, I.; Bokias, G. Langmuir 2005, 21, 9038-9043.
(33) Tsolakis, P.; Bokias, G. Macromolecules 2006, 39, 393-398.
(34) Staikos, G.; Karayanni, K.; Mylonas, Y. Macromol. Chem. Phys. 1997,
198, 2905-2915.
Langmuir, Vol. 23, No. 22, 2007 11253
Scheme 1. Schematic Depiction of the Graft Copolymer
P(AA-co-AMPSA)-g-PDMAM (G48)
ethylamino)propyl)-3-ethyl-carbodiimide hydrochloride (EDC, Aldrich) were used for the synthesis of the graft copolymers.
For the adjustment of the pH citric acid (CA) (Merck) was used.
Water was purified by means of a Seralpur Pro 90C apparatus
combined with a USF Elga laboratory unit. For the SANS
experiments, deuterium oxide (Aldrich) was used.
Polymer Synthesis and Characterization. Amine-terminated
PDMAM was synthesized by free radical polymerization of DMAM
in water at 30 °C for 6 h using the redox couple APS and AET as
initiator and chain-transfer agent, respectively. The polymer was
purified by dialysis against water through the same membrane above
and finally obtained by freeze-drying. Its number-average molecular
weight was determined by end group titration with NaOH after
neutralization with an excess of HCl, using a Metrohm automatic
titrator (model 751 GPD Titrino) and 17000 g/mol was obtained.
A copolymer of AA and AMPSA, P(AA-co-AMPSA), was
prepared by free radical copolymerization of the two monomers in
water, after partial neutralization (90 mol %) with NaOH at pH ≈
6-7, at 30 °C for 6 h, using the redox couple APS/KBS. The product
obtained was then fully neutralized (pH ) 11) with an excess of
NaOH, purified by dialysis against water, and received in its sodium
salt form after freeze-drying. Its composition, determined by acidbase titration and elemental analysis, was 18% in AA units. Its
apparent weight-average molecular weight, Mw ) 2.7 × 105 g/mol,
was determined by static light scattering in 0.1 M NaCl.
The graft copolymer, P(AA-co-AMPSA)-g-PDMAM, was synthesized by a coupling reaction between P(AA-co-AMPSA) and
amine-terminated PDMAM. The two polymers were dissolved in
water at a 1:1 weight ratio. Then, an excess of the coupling agent,
EDC, was added and the solution was stirred for 6 h at room
temperature. Addition of EDC was repeated a second time. The
product was purified with a Pellicon system, equipped with a
tangential flow filter membrane (Millipore, cutoff ) 100 kDa), and
freeze-dried. Its composition in PDMAM side chains was found to
be equal to 48 wt % (using elemental analysis), corresponding to
about 14 chains per graft copolymer. A schematic depiction of the
graft copolymer is presented in Scheme 1. Its apparent molecular
weight, Mw ) 4.8 × 106 g/mol, was determined by static light
scattering in 0.1 M NaCl.
Rheology. Steady-state shear viscosity measurements of semidilute
aqueous polymer mixtures were performed using a Rheometrics SR
200 controlled-stress rheometer, equipped with a cone and plate
geometry (diameter ) 25 mm, angle ) 5.7°, truncation ) 56 µm).
An Anton Paar AMVn automated microviscometer, equipped with
a 1.8 mm diameter glass capillary and a 1.5 mm diameter steel ball,
was used to measure the viscosity of the dilute solutions. The
temperature was fixed at 25 ( 0.1 °C.
Dynamic Light Scattering (DLS). The intensity time correlation
functions g(2)(t) of the polarized light scattering were measured at
θ ) 90° at 24 °C with a full multiple tau digital correlator (ALV5000/FAST) with 280 channels. The excitation light source was a
He-Ne laser (Melles-Griot) operating at 632.8 nm, with a stabilized
power of 17 mW. The incident beam was polarized vertically with
respect to the scattering plane using a Glan polarizer. The scattered
light from the sample was collected through a Glan-Thomson
polarizer (Halle, Berlin) with an extinction coefficient better than
10-7. The samples used were dust-free and optically homogeneous.
The intensity time correlation functions g(2)(t) were analyzed using
the inverse Laplace transformation (ILT) method with the aid of the
CONTIN code.35 From the relaxation times obtained by the ILT
11254 Langmuir, Vol. 23, No. 22, 2007
Sotiropoulou et al.
analysis, the translational diffusion coefficient, DT, of the particles
was determined by means of the equation
DT ) (τq2)-1
(1)
where τ is the relaxation time and q the wave vector given by q )
2πn sin(θ/2)/λ, where n is the refractive index of the medium and
λ the wavelength of the light beam. DT was related to the
hydrodynamic radius, RH, of the particles through the Stokes-Einstein
equation,
RH ) KBT/6πη0D0
(2)
where KB is the Boltzmann constant, T the absolute temperature, η0
the viscosity of the solvent, and D0 the translational diffusion
coefficient at zero concentration.
Static Light Scattering (SLS). SLS measurements were conducted by means of a Model MM1 SM 200 spectrometer (Amtec,
France). An He-Ne 10 mW laser operating at 633 nm was used as
a light source and a complete series of measurements at different
angles and concentrations were conducted for molecular weight
determination. The solutions used, dust-free and optically transparent,
were centrifuged for 2 h at 15.000 turns/min. The refractive index
increment, dn/dC, value was measured by means of a Chromatix
KMX 16 differential refractometer operating also at 633 nm. The
results obtained were subjected to a Zimm analysis, and the molecular
weight of the complexes was determined as the average of the values
found by extrapolation to zero angle and zero concentration.
Small-Angle Neutron Scattering (SANS). SANS measurements
were carried out at the Laboratoire Léon Brillouin (Saclay, France).
The data were collected on beam line PACE at three configurations
(6 Å, sample-to-detector distances 1 m; 7 and 18 Å, 4.55 m), covering
a broad q range from 0.0023 to 0.32 Å-1. Five millimeter light path
quartz cells were used. Empty cell scattering was subtracted and the
detector was calibrated with 1 mm H2O scattering. All measurements
were carried out at room temperature. Data were converted to absolute
intensity through a direct beam measurement, and the incoherent
background was determined with H2O/D2O mixtures.
Atomic Force Microscopy (AFM). AFM images were collected
under ambient conditions (23 °C, 50% RH) using tapping mode
AFM (TM-AFM) on a NanoScope III multimode scanning probe
microscope (Veeco, USA). Silicon tips with a spring constant of 42
N m-1 and a resonance frequency of approximately 320 kHz were
used. In tapping mode,36 the cantilever oscillates at its resonance
frequency (typically 200-400 Hz in air) so that the tip interacts very
briefly with the surface during each oscillation cycle with a small
amplitude (A ∼ 10 nm). The reduction of the cantilever oscillation
from its set point value, due to interactions between the AFM tip
and the sample during the scan, is used to determine the topography
of the surface. To minimize the forces of interaction, the ratio of the
set point value to the free amplitude of the cantilever was maintained
at approximately 0.9 (“light tapping”) by adjusting the vertical
position of the sample. Images were recorded with a resolution of
512 × 512 pixels and a scan rate of 0.5-0.8 Hz. Height and lateral
dimensions of the particles were measured using the Nanoscope
image analysis software (NanoScope V6.13).
Samples were prepared by depositing a drop (5 µL) of polymer
mixture solution (3.6 × 10-5 g‚cm-3) on freshly cleaved mica. The
sample was then gently allowed to evaporate under ambient conditions
in a Petri dish and observed after 20 min.
Preparation of the Polymer Mixture Solutions. Stock solutions
of the mixture G48/PAA were prepared by mixing a 5.50 × 10-2
g/cm3 G48 solution with a 2.10 × 10-2 g/cm3 PAA solution in D2O
for the SANS measurements and a 5.2 × 10-2 g/ cm3 G48 solution
with a 2.00 × 10-2 g/cm3 PAA solution in H2O for all the other
measurements, at pH ) 2.0, adjusted with CA. The mixtures prepared
were considered to contain PAA chains in equivalent quantities with
(35) (a) Provencher S. W. Comput. Phys. Commun. 1982, 27, 213-227. (b)
Provencher S. W. Comput. Phys. Commun. 1982, 27, 229-242.
(36) Zhong, Q.; Innis, D.; Kjoller, K.; Elings, V. B. Surf. Sci. Lett. 1993, 290,
688-692.
Figure 1. Viscosity, η, versus concentration, c, for the polymer
mixture G48/PAA in aqueous solution at pH ) 2.0.
the PDMAM side chains of the graft copolymer G48, that is, in a
unit mole ratio PAA/PDMAM of 1.1/1, according to our previous
study.23 At this point, we consider it useful to point out that a simple
complex should be comprised by one G48 macromolecule and almost
two PAA chains so that its molecular weight should be of the order
of 7 × 105 g/mol. All dilutions were realized with 0.05 M CA (pH
) 2.0). The solutions after their preparation were agitated for 24 h
at room temperature.
Results and Discussion
Rheology. Figure 1 shows the variation of the Newtonian
viscosity, η, vs the concentration, c, for a G48/PAA mixture in
an aqueous solution, at pH ) 2.0. We have chosen the
stoichiometric composition, corresponding to a unit mole PAA/
PDMAM ratio of 1.1, as determined from SANS measurements
at different polymer mixture ratios in a previous work.23 We see
that a critical concentration c* ) 7.0 × 10-3 g/cm3 appears,
separating the dilute from the semidilute concentration regions.
In the semidilute concentration region η increases with c and
follows a scaling law with an exponent equal to 7.5. This high
value shows that a physical gel is formed, due to the interconnection of the anionic backbone chains of G48 in a transient
network, through the hydrogen-bonding interpolymer complexes
formed between its PDMAM side chains and PAA.
DLS. Figure 2 shows the intensity time correlation functions,
g(2)(t), and the ILT distributions for the same as above G48/PAA
mixture in solution at pH ) 2.0 at six different concentrations,
from 0.55 × 10-3 to 1.8 × 10-2 g/cm3. The time correlation
functions curves obtained are generally indicative of a system
comprised of colloidal particles. We observe that, at low
concentrations, c ) 0.55 × 10-3, 1.10 × 10-3, and 1.7 × 10-3
g/cm3, Figures 2a, 2b, and 2c, respectively, single ILT distributions, around 1 × 10-3 s appear. At a higher concentration, c )
6.8 × 10-3 g/cm3, Figure 2d, i.e., close to c*, a broadening of
the distribution to higher times appears, indicating a slowing of
the diffusion times, explained by an increase in the interactions
between the particles. This behavior, which is in accordance
with the viscosity behavior observed above, is even more
accentuated as concentration increases further at c ) 9.0 × 10-3
g/cm3, Figure 2e, that is, higher than c*, where a second peak
in the distribution curve appears at about 1 order of magnitude
higher. Finally, at c ) 1.80 × 10-2 g/cm3, Figure 2f, a third peak
appears at much higher time, while the correlation function curve
is not anymore indicative of any independent particles in the
system. A dramatic slowdown in motion occurs due to the
formation of a transient network taking place in this highconcentration region.
Langmuir, Vol. 23, No. 22, 2007 11255
Figure 2. Intensity time correlation functions, g(2)(t), and ILT distributions for the polymer mixture G48/PAA in aqueous solution at pH
) 2.0, at different concentrations: (a) c ) 0.55 × 10-3 g/cm3; (b) c ) 1.1 × 10-3 g/cm3; (c) c ) 2.2 × 10-3 g/cm3; (d) c ) 6.8 × 10-3
g/cm3; (e) c ) 9.0 × 10-3 g/cm3; (f) c ) 1.8 × 10-2 g/cm3.
Figure 3 shows the concentration dependence of the diffusion
coefficient, D, calculated by means of eq 1, for the three dilute
solutions shown in Figures 2a, 2b, and 2c. The relaxation time
for each solution was obtained by the peak of the corresponding
ILT distribution curve. From the value obtained by extrapolation
to zero concentration, D0 ) 2.33 × 10-8 cm2 s-1, and with use
of eq 2, where we have put T ) 277 K as the room temperature
and η0 ) 0.89 cp for the viscosity of water, a value equal to 105
nm was obtained for the hydrodynamic radius RH of the particles
at infinite dilution.
SLS. SLS measurements at different angles were performed
with dilute solutions and by extrapolation to zero angle and zero
concentration, according to a Zimm plot shown in Figure 4, the
weight average molecular weight, Mw, and the radius of gyration,
RG, of the colloidal nanoparticles were determined. A value of
Mw ) 5.7 × 106 g/mol was obtained for the molecular weight,
somewhat higher but comparable to the value M ) 4.5 × 106
g/mol, calculated after SANS measurements in the following.
An aggregation number equal to 6-8 can be calculated on the
basis of these molecular weight results, showing that each colloidal
nanoparticle should be comprised of 6-8 graft copolymer chains
and 12-16 PAA chains. Regarding the radius of gyration, the
value RG ) 85 nm was obtained, which combined with the
hydrodynamic radius, RH ) 105 nm, found above shows that the
colloidal particles should be of a spherical form, as RG/RH is
close to the square root of 3/5.
SANS. Figure 5 shows the variation of the SANS intensity,
I, versus the scattering wave vector, q, for the same G48/PAA90
mixture at six different concentrations in D2O, at pH ) 2. The
scattered intensity can be discussed separately for three different
q regions. At low q, roughly q < 0.01 Å-1, typical Guinier
scattering is found at low concentration, indicative of the finite
11256 Langmuir, Vol. 23, No. 22, 2007
Sotiropoulou et al.
q0 ) 0.0108 Å-1, and φ ) 1.85 × 10-2, calculated by taking
into account only the compact complex particles formed between
the PDMAM side chains of G48 and the PAA chains and their
mass density, dc ) 1.28 g/cm3, determined elsewhere,23 Rdry
becomes equal to 9.5 nm.
The low q intensity region corresponds to a first approximation
to the Guinier regime of the scattering of individual noninteracting,
finite-sized objects;38 their radius leads to a characteristic decrease
in I, whose magnitude is related to their mass. If the objects are
spheres of radius Rc,
I ) I0 exp(-Rc2q2/5)
(4a)
I0 ) φ∆F2V0
(4b)
with
Figure 3. Variation of the translational diffusion coefficient DT as
a function of the concentration, c, in the low-concentration region,
for the mixture G48/PAA at pH ) 2.0, and extrapolation to zero
concentration.
Figure 4. Zimm plot for the polymer mixture G48/PAA at pH )
2.0.
size of aggregates. At intermediate q, 0.01 < q < 0.1 Å-1, we
observe that I decreases abruptly, following a scaling law of the
form I ∼ q-d. As the values of the exponent d vary between 3.5
and 4.0, the presence of three-dimensional objects with smooth
or fractal surfaces is indicated,37 which we attribute to the insoluble
hydrogen-bonding interpolymer complexes formed between the
PDMAM side chains of the graft copolymer and the PAA
chains.17,23 At high q, q > 0.1 Å-1, finally, chain scattering is
found, which should be attributed to the anionic backbone of the
graft copolymer, comprising the hydrophilic shell of the colloidal
particles formed.
Furthermore, as the concentration becomes higher than c*,
the critical overlapping concentration, the form of the intensity
curves changes, with a tendency to shift to lower values at low
q and to exhibit a structural peak, at around 0.01 Å-1, which
becomes clear only in the most concentrated solution, Figure 3f,
reflecting the interactions between the objects. By considering
that it corresponds to the most probable distance between them,
we can apply a cubic lattice model based on the mass conservation
of the complex particles, with the distance between the particles
given through D ) 2π/q0. Since the volume, V, of each particle
can be estimated by V ) φD3, where φ is the volume fraction
of the particles, their “dry” radius, Rdry, can be calculated by
3
Rdry )
x
6π2φ
q03
(3)
In the case of the most concentrated solution, Figure 5f, where
(37) Higgins, J. S.; Benoı̂t H. C. In Polymers and Neutron Scattering; Oxford
Science Publications: Clarendon Press: Oxford, 1994.
where V0 denotes the dry volume of an individual object, φ the
volume fraction of the objects, and ∆F the scattering contrast
between the solvent and the dry polymer.
Then from eq 4b, by taking ∆F ) 5.0 × 1010 cm-2,23 we obtain
the value of 165 cm-1 for the intensity at zero q, I0. Using this
value in the Guinier form expressed by eq 4a, we obtain a relatively
good fit for the data of Figure 5f, if we use a value equal to 16
nm for the radius, Rc, of the compact particles. We also observe
that we have relatively good Guinier fitting for all the concentrations measured by using as I0 the value occurring from the initially
estimated quantity for the most concentrated solution, 165 cm-1,
adjusted each time proportionally to the concentration. The value
for the radius of the particles obtained is practically stable at
16-17 nm. It should also be considered as a “wet” radius
representing about 80% hydrated particles. Moreover, it should
be compared to the hydrodynamic radius of the particles, RH )
105 nm, obtained from DLS measurements in dilute solution. It
is noteworthy that this hydrodynamic radius includes not only
the insoluble core of the compact hydrogen-bonding interpolymer
complexes formed between PAA and the PDMAM side chains
of the graft copolymer but also a hydrophilic shell comprised of
its anionic backbone. A representative schematic depiction of
the colloidal nanoparticles formed is presented by Scheme 2.
The hydrophilic shell is comprised of loops and single strands
of the anionic backbone, extended, due to their charge and the
low ionic strength of the solution, while their length should be
related to the distribution of the PDMAM side chains in the ionic
backbone and its length estimated to be over the 330 nm based
on its molecular weight.
From the volume of the particle we can also obtain the
molecular mass, Mc, of the dry complex particle
Mc ) VdcNA
(5)
where NA is Avogadro’s number. Equation 5 gives Mc ) 2.8 ×
106 Da, corresponding to a value equal to 4.5 × 106 Da for the
whole particle. This core molecular weight value also implies
that each particle contains about 90 PDMAM side chains involving
more than six graft copolymer chains. This leads to the formation
of a transient network, explaining the increase in viscosity
observed in semidilute solution, Figure 1, and the gel formation
already studied.17,21
AFM. Figure 6 represents an AFM image showing globular
particles homogeneously distributed on the mica surface. The
particles obtained are characterized by lateral dimensions of about
(38) Lindner, P., Zemb, Th., Eds.; Neutrons, X-rays and Light: Scattering
Methods Applied to Soft Matter; North-Holland, Delta Series; Elsevier: Amsterdam,
2002.
Langmuir, Vol. 23, No. 22, 2007 11257
Figure 5. SANS intensity variation vs the wave vector q, for the G48/PAA polymer mixture in solution in D2O, at pH ) 2, at different
concentrations: (a) c ) 1.6 × 10-3 g/cm3; (b) c ) 3.2 × 10-3 g/cm3; (c) c ) 6.3 × 10-3 g/cm3; (d) c ) 9.5 × 10-3 g/cm3; (e) c ) 1.9
× 10-2 g/cm3; (f) c ) 3.8 × 10-2 g/cm3.
60 nm, with a height of about 1.5-2 nm. It is well-established
that lateral dimensions, determined by AFM, are overestimated
due to the convolution effect of the AFM tip when scanning
small objects.39 Assuming a tip radius of 10 nm (the actual size
of commercial AFM tips is given between 5 and 15 nm), we can
estimate a true lateral size around 22.5 nm. From these dimensions
and with the assumption of spherical cap geometry, the particles
volume deposited on the substrate is about 3200 nm3. It appears
that this estimated volume is lower than the one of the insoluble
core in solution, as determined by SANS (17150 nm3), but it is
in a fairly good agreement with the dry volume of the core, as
it has been estimated to be 80% hydrated. On the other hand,
(39) Westra, K.L.; Mitchell, A.W.; Thomson, D.J. J. Appl. Phys. 1993, 74,
3608-3610.
Scheme 2. Negatively Charged Colloidal Particles Formed
through Hydrogen-Bonding Interpolymer Complexation of
PAA with the PDMAM Side Chains of the Graft Copolymer
P(AA-co-AMPSA)-g-PDMAM (G48), at Low pH
11258 Langmuir, Vol. 23, No. 22, 2007
Sotiropoulou et al.
Figure 6. Tapping mode AFM picture of the G48/PAA polymer mixture deposited on mica and observed under ambient conditions.
the drying procedure used in the preparation of the sample before
AFM imaging, which is useful to immobilize the polymer on the
mica substrate (both mica and the polymer are negatively charged),
may induce a collapse of the colloidal nanoparticles due to a
conformation change upon deposition on the mica and/or a
possible dehydration of the polymer. As a consequence, the
observed dimension and shape of the nanoparticles observed by
AFM should look like the one of the dry core. This finding
emphasizes the fact that it is a multiscale organized particle with
a central hydrophobic core, hydrated up to 80%, comprised of
the hydrogen-bonding IPC formed between PAA and the
PDMAM side chains of G48, and a hydrophilic shell made of
the P(AA-co-AMPSA) anionic backbone of the G48.
Conclusions
We have studied the hydrogen-bonding interpolymer complexation between PAA and PDMAM grafted onto a negatively
charged backbone (P(AA-co-AMPSA)) by viscometry, dynamic
and static light scattering, SANS, and AFM. The results obtained
in aqueous solution at pH ) 2.0 revealed a structured system
consisting of anionic colloidal nanoparticles. According to
dynamic and static light scattering results, spherical particles are
formed with a hydrodynamic radius of about 105 nm. They are
comprised of a compact core of PAA/PDMAM hydrogen-bonding
interpolymer complexes and a hydrophilic shell of anionic P(AAco-AMPSA) chains. SANS measurements showed that the
hydrophobic core presents a radius of 16-17 nm and a molar
mass of 2.8 × 106 Da. AFM revealed the formation of particles
with a size approaching that of the hydrophobic core.
Acknowledgment. This research project has been supported
by the European Commission under the 6th Framework Programme through the Key Action: Strengthening the European
Research Area, Research Infrastructures. Contract No. HII3CT-2003-505925.
LA701561Y
Linear and nonlinear viscoelastic behavior of very
concentrated plate-like kaolin suspensions
Frédéric Bossarda)
Laboratoire de Rhéologie UMR 5520, BP 53, Université Joseph Fourier,
1301 rue de la Piscine, 38041 Grenoble Cedex 9, France
Michel Moan and Thierry Aubry
Laboratoire de Rhéologie, 6 avenue Le Gorgeu - CS93837,
29238 Brest Cedex 3, France
(Received 8 December 2006; final revision received 1 August 2007兲
Synopsis
The viscoelastic behavior of very concentrated and electrostatically stabilized suspensions of
kaolinite particles has been investigated in the linear and nonlinear regime as a function of volume
fraction, ionic strength and in the presence of polymer at various concentrations. Material
properties such as linear viscoelastic moduli and cohesive energy density are extensively enhanced
by either increasing volume fraction or decreasing ionic strength. Attention has been paid to the
large amplitude oscillatory shear behavior of concentrated suspensions of plate-like particles,
characterized by a hump in G⬙ curves. Rheological investigation shows the extreme sensitivity of
the intensity of the strain hardening in G⬙ to excluded volume, electrostatic and steric interactions.
A physical interpretation of this nonlinear behavior has been proposed. © 2007 The Society of
Rheology. 关DOI: 10.1122/1.2790023兴
I. INTRODUCTION
The anisometric character of kaolinite particles, combined with the electrostatic properties of their surface, gives these clay suspensions quite specific structural properties.
Indeed, at volume fraction ␾ ⬎ ␾* ⬃ 0.1 corresponding to the onset of excluded volume
interactions, Jogun and Zukoski 共1996, 1999兲 suggest that plate-like kaolinite particles
are aligned within domains. More recently, scanning electron cryomicroscopy observations have pointed out that very concentrated kaolinite suspensions are organized in
quasi-continuous neighboring domains of aligned, close-packed particles exhibiting a
nematic order over few micrometers 关Moan et al. 共2003兲兴. Such microstructural organization confers peculiar flow properties to kaolinite suspensions that have been thoroughly
investigated. In the intermediate shear rate region, the presence of a “hesitation” point in
the steady state flow curve, associated with a negative minimum of the first normal stress
difference, was shown to be reminiscent of nematic liquid crystalline polymer behavior
关Moan et al. 共2003兲兴. The interpretation proposed by the authors suggests a competition
between shear forces that tend to align domains and interactions between neighboring
a兲
Author to whom correspondence should be addressed; electronic mail: [email protected]
J. Rheol. 51共6兲, 1253-1270 November/December 共2007兲
0148-6055/2007/51共6兲/1253/18/$27.00
1253
1254
BOSSARD, MOAN, AND AUBRY
domains that oppose to their mutual alignment, leading to a progressive uniform alignment of the domains in the flow direction, as the shear forces increase.
By considering each plate-like clay particle with its ionic double layer as a new
effective particle, the decrease of the ionic strength is known to increase the effective
particle size through the expansion of the ionic double layer. Concurrently, long-range
electrostatic repulsions are gradually strengthened, favoring the edge/face perpendicular
configuration by minimizing their mutual repulsions 关Mourchid et al. 共1995兲; Meyer et al.
共2001兲兴. Consequently, on decreasing the ionic strength, a competition between the tendency of plate-like particles to align and the tendency to an isotropic ordering has been
shown 关Rowan and Hansen 共2002兲兴. For large plate-like particles, such as bentonite, the
relative change in the aspect ratio is less sensitive to ionic strength variation. So that,
isotropic structure is expected to be favored when decreasing the ionic strength.
In the presence of polymers, three different mechanisms may occur: For nonadsorbing
polymer chains at high concentrations, depletion flocculation takes place through phase
separation 关Sperry et al. 共1981兲兴. For polymer chains that adsorb on clay surface, adsorbing macromolecules may bridge particles at concentrations below the saturation of accessible clay surface 关Lafuma et al. 共1991兲; Spalla and Cabane 共1993兲兴. They may form
a polymeric layer around the particle at concentrations above the saturation of accessible
clay surface, leading to the increase of steric interactions 关Napper 共1983兲; de Gennes
共1987兲兴. Yziquel et al. 共1999a, 1999c兲 have shown a great enhancement of rheological
properties of kaolinite suspensions at ␾ = 0.34 by adding high-molecular-weight polymers, such as carboxymethyl cellulose 共CMC兲 or polyvinyl alcohol 共PVA兲.
Flow curves express the rheological response at large deformations, whereas viscoelastic investigation techniques are measurements performed to provide material properties of structured systems at very low deformations. The usual test consists of measuring the frequency dependence of storage and loss moduli in the linear viscoelastic regime,
i.e., in the low shear strain amplitude domain where the response is sinusoidal and both
G⬘ and G⬙ moduli are independent of strain amplitude. This spectromechanical technique
has mainly two advantages: it is very sensitive to the microstructure and it can be treated
in a rigorous mathematical framework 关Macosko 共1994兲兴. Besides, some works focus on
the rheological response of complex fluids in the nonlinear viscoelastic regime, where G⬘
and G⬙ moduli are dependent on both frequency and strain amplitude. Among these
measurements, the large amplitude oscillatory shear test 共LAOS兲, which consists of measuring viscoelastic moduli as a function of shear strain amplitude from the linear up to the
nonlinear regime at a fixed frequency, has been shown to be useful to investigate the
microstructural state of various complex fluids 关Yosick and Giacomin 共1996兲; Yosick et
al. 共1997兲; Hyun et al. 共2002, 2003, 2006兲; Sim et al. 共2003a, 2003b兲兴. However, LAOS
measurements carried out using commercial rheometers must be considered with caution
since G⬘ and G⬙ moduli lose their physical meaning beyond the linear viscoelastic regime. Indeed, the stress becomes no longer sinusoidal and contributions of higher harmonics to viscoelastic moduli may be non-negligible 关Dealy and Wissbrun 共1990兲兴. The
harmonic contributions arising from nonlinear effects can be analyzed using Fourier
transformation methods 关Wilhelm et al. 共1998, 1999, 2000, 2002兲; See 共2001兲; Karis et
al. 共2002a, 2002b兲兴 or using graphical analysis by drawing a Lissajous curve 关Hyun et al.
共2003兲兴. Very recently, a descriptive approach of viscoelasticity, extended in the nonlinear
regime, has been proposed to interpret LAOS data 关Cho et al. 共2005兲兴. Indeed LAOS
measurements have been used to classify the viscoelastic behavior of complex fluids in
four categories: type I, strain thinning 共G⬘ and G⬙ decrease兲; type II, strain hardening 共G⬘
and G⬙ increase兲; type III, weak strain overshoot 共G⬘ decrease and hump in G⬙ curve兲;
type IV, strong strain overshoot 共hump in both G⬘ and G⬙ curves兲关Hyun et al. 共2002兲;
VERY CONCENTRATED KAOLIN SUSPENSIONS
1255
Sim et al. 共2003a兲兴. The LAOS behavior of concentrated kaolinite suspensions has shown
to be classified in type III but the microscopic origin of G⬙ hump has been scarcely
discussed up to now. Simulation analysis using a modified Jeffreys model with a single
relaxation time and an energy dependent kinetic equation has shown to fit correctly the
nonlinear behavior of kaolinite suspension 关Yziquel et al. 共1999a兲兴. This model assumes
that the breakdown of the suspension microstructure is related to the rate of energy
dissipated by oscillatory shear. Studies performed on fumed silica suspensions have
shown that a particle network is required to develop the strain hardening in G⬙ 关Yziquel
et al. 共1999b兲兴. According to the authors, the nonlinear viscoelastic behavior of concentrated suspensions is governed by microstructural changes, which result from the competition between the breakup of the network under flow and its buildup due to Brownian
motion. The dissipative energy per unit volume has been shown to depend on particle
size, nature of the surface and suspending medium, volume fraction and shear strain
amplitude. Despite these studies, the physical mechanism of the strain hardening in G⬙ is
still not fully understood.
In this article, we have investigated the influence of volume fraction, ionic strength
and polymer concentration separately, to discriminate the contribution of excluded volume, electrostatic and steric interactions, respectively, to the linear and nonlinear viscoelastic behavior. The main objective of the paper is to give some physical insight into
the microstructural origin of the G⬙ hump.
II. MATERIALS AND EXPERIMENTAL METHODS
A. Kaolinite
The clay particles used in this study are kaolinite particles, commercialized by Engelhard Corporation 共New Jersey兲 and referenced as Miragloss 91. Kaolinite particles consist of alumino silicate layers, responsible for their plate-shaped geometry. The average
diameter of kaolinite particles d ⬃ 032 ␮m, obtained from the particle size distribution
determined with a particle size analyzer 共laser Malvern Mastersizer 2000兲, is consistent
with the diameter value of 0.30 ␮m given by the supplier; the aspect ratio d / h is about
10, h being the thickness of the plate.
The charge on the faces of the kaolin particles is negative and pH independent, while
the charge on the edges changes from positive to negative values with increasing pH, due
to the coexistence of both positive and negative charges. Between pH 5 and 9, the edge
surface is expected to be negatively charged 关Lee et al. 共1991兲兴. The kaolin powder was
dispersed in an aqueous sodium phosphate buffer solution with an ionic strength I ⬃ 4
⫻ 10−3 M. A washing process, consisting of a centrifugation and redispersion sequence,
was repeated until the pH 共⬃7.3兲 and the ionic strength of the suspending medium were
the same as the original buffer solution 关Jogun and Zukoski 共1996兲兴. As shown by Nicol
and Hunter 共1970兲, phosphate ions preferentially condense on the edges, leading to the
neutralization of possible positive charges. Under these pH and ionic strength conditions,
the negative charge of both the basal planes and the edges of the particles gives rise to
electrostatic repulsive interactions, minimizing particle aggregation. Measurements have
been performed at ionic strength ranging from 3 ⫻ 10−3 M to 1.3⫻ 10−2 M and volume
fraction ␾ from 0.33 to 0.55 in order to study the influence of electrostatic and steric
interactions on viscoelastic behavior. All suspensions have a volume fraction much higher
than the critical volume fraction ␾* ⬃ 0.1, corresponding to the onset of excluded volume
interactions.
1256
B. Polymer
The polymer used is a commercial 共hydroxypropyl兲 guar, 共HPG兲, synthesized by Fratelli Lamberti s.p.a. 共Albizzate, Italy兲. The weight average molecular weight is about 2
⫻ 1016 corresponding to a degree of polymerization of about 3000, and the dispersity
index is close to 1.5. This polysaccharide contains an average of one hydrophilic substituent 共hydroxypropyl or hydroxybutyl group兲 per monomer. The intrinsic viscosity of
HPG macromolecules is ⬃1200 cm3 / g corresponding to a radius of gyration of about
90 mm.
C. Adsorption measurements
Kaolinite suspensions in the presence of polymer have been prepared by dispersing the
proper amount of clay particles into polymer solutions at ionic strength I = 4 ⫻ 10−3 M.
Suspensions are stirred 24 h at room temperature before measurements. Samples containing polymer were first centrifuged at 18,000 rpm at 25 ° C during 1 h in order to separate
kaolinite particle covered by polymer chains, located at the bottom, from the nonadsorbed
chains located in the supernatant. A direct concentration measurement by the total organic
carbon technique was used to determine the equilibrium concentration of the free chain
concentration Cequ in the supernatant. For this purpose, a small sample of supernatant was
heated at 68 ° C in an O2 atmosphere, so that all carbon atoms are to be found in CO2,
whose concentration was determined by spectrometry; the free polymer concentration
was inferred from knowledge of the chemical structure of HPG.
The polymer-adsorbed amount ⌫ was calculated from the difference between the initial polymer concentration and the equilibrium concentration.
D. Rheometry
Viscoelastic measurements were performed on a controlled strain rheometer 共ARES,
Rheometric Scientific兲 equipped with either a cone-and-plate geometry 共diameter⫽
50 mm, cone angle⫽0.04 rad兲 or a parallel-plate geometry 共diameter⫽25 mm, gap⫽
2 mm兲. The absence of a significant slip at the wall was verified by varying the gap from
0.2 to 2 mm. Viscoelastic measurements are influenced by the stress history imposed on
the material before experiment, for example, during the loading of the sample in the
rheometer. So, a protocol was defined to assure the desired reproducibility: after loading
in the rheometer, the sample is kept at rest for a fixed time before starting the rheological
measurement. The resting time has been preliminary, determined from oscillatory measurements performed in the linear viscoelastic domain: immediately after loading the
sample, the time evolution of the elastic modulus G⬘ at a frequency of 1 Hz is followed
and the time needed to attain a G⬘ constant value is determined. This time, which is about
5 min, can be considered as the time needed to recover an equilibrium structure. Moreover, we have systematically verified that the nonzero normal force, which appears during the loading of the sample, has decayed to zero before starting a test. A thin layer of
low viscosity silicone oil was spread over the air/suspension interphase in order to prevent solvent evaporation. The presence of edge instabilities, such as edge fracture 关Larson
共1992兲兴, has never been detected by visual observation of the air/suspension interface
during measurement. Compared to other systems, such as polymer solutions or melts
which have large normal stresses, the absence of visible edge instabilities is probably due
to the low elasticity level of the suspensions studied. At all ionic strengths, volume
fractions and polymer concentrations investigated, electrostatic repulsive interactions,
combined with excluded volume interactions are strong enough to allow a good disper-
1257
FIG. 1. Storage modulus G⬘ and loss modulus G⬙, at frequency of 1 Hz, as a function of shear strain amplitude,
␾ = 0.55.
sion and stability of the suspension over a long period of time 共a few days兲 and prevent
the samples from sedimentation during rheological tests 关Bossard 共2001兲兴. All tests were
performed at 25 ° C.
III. RESULTS AND DISCUSSION
A. Influence of volume fraction
Figure 1 shows the shear strain amplitude dependence of the storage modulus and the
loss modulus measured at the frequency of 1 Hz for a suspension at ␾ = 0.55. The viscoelastic behavior is representative of those obtained for any suspensions tested. As shear
strain amplitude increases, both moduli exhibit a constant value G0⬘ and G0⬙, with G0⬘
⬎ G0⬙, until a critical shear strain amplitude ␥c, which defines the extent of the linear
viscoelastic regime. Above ␥c, G⬘ modulus decreases gradually with increasing shear
strain amplitude while G⬙ modulus exhibits a strain hardening characterized by a hump
with a maximum value, noted Gmax
⬙ , for a shear strain amplitude ␥max.
Let us consider first the linear viscoelastic regime. Figures 2共a兲 and 2共b兲 show G0⬘ and
G0⬙ moduli and critical shear strain amplitude ␥c as a function of volume fraction, respectively. In the narrow volume fraction range investigated, both moduli in the linear regime
increase sharply with ␾. The volume fraction dependence of G0⬘ and G0⬙ moduli is properly described by a power law, with an exponent of about 17 and 11, respectively. Similar
volume fraction dependence of G0⬘, with an exponent lying from 10 to 20, has been
mentioned for sterically interacting hard sphere suspensions 关Rabaioli et al. 共1993兲; Rao
et al. 共2006兲兴. Due to the anisometric character of kaolinite particles, volume fraction
effects will be discussed in terms of excluded volume interactions rather than steric
interactions. Such volume fraction dependence of the material parameters G0⬘ and G0⬙
suggests that excluded volume interactions play an important role in the elasticity and
viscosity enhancement of kaolinite suspensions.
The linear viscoelastic behavior can be analyzed in terms of cohesive energy density
Ec corresponding to the work needed to break the structure
1258
FIG. 2. 共a兲 Storage modulus G⬘0 and loss modulus G⬙0 in the linear regime; 共b兲 critical shear strain amplitude ␥c
and 共c兲 cohesive energy density Ec as a function of volume fraction, ␻ = 1 Hz.
Ec =
冕
␥c
␶d␥ .
共1兲
0
Since the shear stress in the linear viscoelastic regime is given by
␶ = G0⬘␥ ,
the cohesive energy density is then
共2兲
1259
FIG. 3. Reduced storage modulus G⬘ / G⬘0 as a function of reduced shear strain amplitude ␥0 / ␥c at volume
fractions ␾ = 0.40 共〫兲, 0.45 共䊏兲, 0.50 共䉭兲, and 0.55 共쎲兲, ␻ = 1 Hz. Inset: reduced storage modulus G⬘ / G⬘0 of
the suspension at ␾ = 0.55 as a function of reduced shear strain amplitude ␥0 / ␥c at different frequencies.
1
Ec = ␥2c G0⬘ .
2
共3兲
As depicted in Fig. 2共c兲, the cohesive energy density of kaolinite suspensions scales as
␾7. Such power law dependence of Ec has been observed for montmorillonite 关Sohm and
Tadros 共1989兲; Aubry and Moan 共1997兲兴 and Laponite 关Ramsay 共1986兲兴 suspensions with
smaller exponents, 3 and 1.8, respectively. From a phenomenological point of view, the
increase of the cohesive energy density with increasing volume fraction is a direct consequence of the enhancement of excluded volume interactions between plate-like particles, due to the reduction of the average inter-particle distance.
Let us consider now the nonlinear viscoelastic regime. In order to compare qualitative
effects of volume fraction on viscoelastic behavior, G⬘ and G⬙ moduli have been normalized by their respective values in the linear regime G0⬘ and G0⬙. Figure 3 represents the
reduced storage modulus G⬘ / G0⬘ as a function of reduced shear strain amplitude ␥0 / ␥c at
various volume fractions. All reduced storage moduli can be plotted on a master curve.
To have a rheo-physical insight into the viscoelastic behavior of concentrated suspensions, measurements have been carried out at various frequencies. As shown in the inset
for a suspension at ␾ = 0.55, the elastic behavior is frequency independent. These results
point out that the elastic response of kaolinite suspensions is not related to the time scale
of the mechanical stress but depends mainly on the shear strain amplitude. Figure 4
shows the reduced loss modulus G⬙ / G0⬙ as a function of shear strain amplitude ␥0 at
different volume fractions. At ␾ = 0.40, the loss modulus decreases with increasing shear
strain amplitude whereas the strain hardening in G⬙ appears at intermediate shear strain
amplitudes for ␾ ⬎ 0.40 and its intensity increases with ␾, i.e., with decreasing interparticle distance. As observed for the storage modulus, the inset shows that the loss
1260
FIG. 4. Reduced loss modulus G⬙ / G⬙0 as a function of shear strain amplitude ␥0 at volume fractions ␾ = 0.40
共〫兲, 0.45 共䊏兲, 0.50 共䉭兲, and 0.55 共쎲兲, ␻ = 1 Hz. Inset: reduced loss modulus G⬙ / G⬙0 of the suspension at ␾
= 0.55 as a function of reduced shear strain amplitude ␥0 / ␥c at different frequencies.
modulus profile is frequency independent. It has to be noticed that such strain hardening
in G⬙ has been already observed for kaolinite 关Jogun and Zukoski 共1999, 1996兲; Yziquel
et al. 共1999a, 1999c兲兴, Laponite 关Avery and Ramsay 共1986兲兴 and montmorillonite suspensions 关Marchal et al. 共1996兲兴. However, this nonlinear G⬙ feature at intermediate shear
strain amplitudes has been scarcely discussed and interpreted 关Avery and Ramsay
共1986兲兴. The strain hardening in G⬙ is expected to depend on inter-particle interactions,
that is, on inter-particle distance. As volume fraction lies from 0.4 to 0.55, the interparticle distance, noted e and estimated from a packed organization of piled clay particles
using Eq. 共4兲, ranges approximately from 50 to 20 nm, which is small compared to the
average plate diameter d ⬃ 320 nm.
e=h
冉冊
1−␾
.
␾
共4兲
At such dense packing, an oscillatory shear induces complex relative motions of two
very close neighboring particles. The relative displacement has been evaluated, considering the simple case of two parallel particles with an inter-particle distance e, oriented
with an angle ␣ relative to the velocity field, as presented in Fig. 5. For this purpose, the
displacement gradient tensor 共ⵜ␰兲 has been determined using Eq. 共5兲. Details are reported
in the Appendix:
共ⵜ␰兲 =
冢
␥0
sin 2␣ ␥0 cos2 ␣
2
␥0
␥0 sin2 ␣
sin 2␣
2
−
冣
.
共5兲
1261
FIG. 5. Schematic illustration of two neighboring plate-like particles under oscillatory shear.
Thus the relative displacements along directions parallel and orthogonal to the basal
plane are given by the following relations:
⌬d
e
= ␥0 cos2 ␣ ,
d
d
共6兲
⌬e ␥0
=
sin 2␣ .
2
e
共7兲
These equations show that kaolinite particles have a complex motion: the displacement, ⌬d, along the basal plane corresponds to a sliding motion and the displacement,
⌬e, along the orthogonal direction to the basal plane, corresponds to alternative pushing
and pulling apart basal surfaces of neighboring particles.
For ␣ = 0, the relative displacement is a slipping motion only and for ␣ = 45°, ⌬e is
maximum. These two relative displacements have been estimated for clay suspensions at
␾ = 0.55 in the arbitrary case ␣ = 30° at shear strain amplitude ␥c = 2% and ␥max = 30%. At
shear strain amplitude ␥c, ⌬d / d and ⌬e / e are small: ⬃0.1% and 0.6%, respectively.
Therefore, under shear strain amplitude ␥0 艋 ␥c, the relative displacement of clay particles is too weak to disrupt the microstructure. As shear strain amplitude increases,
relative displacements ⌬d / d and ⌬e / e increase and reach 1.6% and 9.6%, respectively, at
shear strain amplitude ␥max = 30%. Thus, along the basal plane, displacement ⌬d of kaolinite plates is small compared to their average diameter but displacement ⌬e is nonnegligible compared to the inter-particle distance. To get a physical insight into the origin
of the strain hardening in G⬙, let us consider the repulsive electrostatic interactions and
excluded volume interactions. The fluctuation of the inter-particle distance induced by
oscillatory shear measurement is responsible for the variation of both electro-viscous
dissipation, through the deformation of the diffuse charge clouds surrounding particles,
and steric interactions between neighboring particles. The two contributions are combined: excluded volume interactions force close-packed plate-like particles to align,
which enhances electrostatic interactions. On increasing volume fraction, inter-particle
distance decreases, leading to strengthen the electro-viscous dissipation and steric interactions. Consequently, the increase of the strain hardening intensity in G⬙ with increasing
␾ originates from the increase of both excluded volume and electrostatic interactions,
governed by the fluctuation of the inter-particle distance during oscillatory shear measurement. This physical interpretation will be confirmed in the following section devoted
to the influence of ionic strength.
1262
FIG. 6. 共a兲 Storage modulus in the linear regime, 共b兲 critical shear strain amplitude ␥0 and 共c兲 cohesive energy
density Ec as a function of ionic strength at volume fractions ␾ = 0.40 共〫兲, 0.45 共䊏兲, 0.50 共䉭兲, and 0.55 共쎲兲,
␻ = 1 Hz.
B. Influence of ionic strength
Figures 6共a兲 and 6共b兲 show the plateau storage modulus G0⬘ and the critical shear strain
amplitude ␥c, respectively, measured at ␻ = 1 Hz, as a function of ionic strength, at
volume fractions ␾ = 0.40, 0.45, 0.50, and 0.55. For all volume fractions investigated, the
plateau storage modulus decreases and the critical shear strain amplitude increases with
increasing ionic strength. The ionic strength dependence of G0⬘ and ␥c induces a decrease
of the cohesive energy density with increasing ionic strength, as depicted by Fig. 6共c兲. It
has to be noticed that the decay of Ec is more and more pronounced as volume fraction
1263
FIG. 7. 共a兲 Reduced storage modulus G⬘ / G⬘0, 共b兲 reduced loss modulus G⬙ / G⬙0 as a function of reduced shear
strain amplitude ␥0 / ␥c at ionic strengths I = 3 ⫻ 10−3 M 共䊏兲, 4 ⫻ 10−3 M 共䉭兲, 6.8⫻ 10−3 M 共쎲兲, 10−2 M 共〫兲
and 1.3⫻ 10−2 M 共䉲兲, ␾ = 0.55 and ␻ = 1 Hz.
increases. At ionic strength I = 4 ⫻ 10−3 M, the Debye length ␬−1 is close to 4 nm, that is
⬃20% of the inter-particle distance e at ␾ = 0.55. Consequently, basal surfaces bearing
negative charges undergo strong repulsive interactions at low ionic strength. As ionic
strength increases, repulsive interactions are gradually screened, leading to a decrease of
the cohesive energy density. Such effect is enhanced when the inter-particle distance
decreases, i.e.,when the volume fraction increases.
The ionic strength also modifies qualitatively the viscoelastic behavior in the nonlinear
regime. Figures 7共a兲 and 7共b兲 show the reduced storage modulus G⬘ / G0⬘ and the reduced
loss modulus G⬙ / G0⬙ as a function of reduced shear strain amplitude ␥0 / ␥c, measured at
␻ = 1 Hz, for a suspension at ␾ = 0.55 at different ionic strengths, from I = 3 ⫻ 10−3 M to
1.3⫻ 10−2 M. The reduced storage modulus G⬘ / G0⬘ is well described by a master curve
for all ionic strengths investigated, whereas the G⬙ / G0⬙ hump is progressively reduced as
ionic strength is increased. As shown in Fig. 8, the decay of the intensity of strain
hardening in G⬙ with increasing ionic strength comes out for all volume fractions inves-
1264
FIG. 8. Reduced intensity of the G⬙ peak as a function of ionic strength, at volume fractions ␾ = 0.40 共〫兲, 0.45
共䊏兲, 0.50 共䉭兲, and 0.55 共쎲兲, ␻ = 1 Hz.
tigated, and it is more pronounced as volume fraction increases. The ionic strength
dependence of the intensity of the strain hardening in G⬙ is attributed to the decrease of
the Debye length ␬−1 that weakens electro-viscous effects.
C. Influence of polymer concentration
In this third section, HPG at various concentrations is added to concentrated suspensions at ␾ = 0.45 and I = 4 ⫻ 10−3 M in order to modify the nature of the inter-particle
interactions that govern the rheological behavior. Preliminary adsorption measurements
have been carried out to define the saturation concentration, corresponding to the total
coverage of kaolinite particles by polymer chains. Measurements that request the centrifugation of suspensions cannot be performed at ␾ = 0.45 for which a sufficient volume
of supernatant cannot be obtained. Consequently, adsorption measurements have been
performed at solid/liquid ratios S / L ranging from 0.46% to 20%, corresponding to volume fractions from 0.0018 to 0.0714, respectively. Figure 9 shows the adsorption isotherm of HPG at various solid/liquid ratios. As equilibrium concentration increases, all
adsorption isotherms are characterized by a sharp increase of the adsorbed amount, followed by an adsorption plateau marking the saturation of kaolinite surfaces. Adsorption
isotherms can be correctly fitted to Langmuir equation
K . Cequ
⌫
=
,
⌫m 1 + K . Cequ
共8兲
where ⌫m is the saturated adsorption and K is the equilibrium constant, in the very dilute
limit. This model has been applied successfully to describe the adsorption isotherm of
HPG on Laponite particles 关Aubry et al. 共2002兲兴. Equation 共8兲 is used as a fit equation to
determine the saturated adsorption at different S / L ratios. Adsorption at saturation ⌫m,
reported as a function of S / L ratios in inset of Fig. 9, decreases upon increasing volume
fraction. Such decay is due to the decrease of accessible particle surfaces for polymer
chains as inter-particle distance decreases 关Lee et al. 共1991兲; Argillier et al. 共1996兲;
Nabzar et al. 共1986兲兴. In the S / L ratio range investigated, the saturated adsorption decay
1265
FIG. 9. Adsorption isotherms of HPG on kaolinite particles at solid/liquid ratio S / L = 0.46% 共䉭兲, 1% 共쎲兲, 2%
共䊐兲, 4% 共䉱兲, 10% 共〫兲, and 20% 共䉲兲. Inset: adsorption at saturation, ⌫m, as a function of S / L.
is correctly fitted by a power law. Assuming that the adsorption mechanism is not modified at high S / L ratios, the adsorption at saturation at S / L = 210% 共corresponding to ␾
= 0.45兲 can be estimated through data extrapolation. The extrapolated ⌫m value is about
1 mg/ g, which corresponds to a polymer concentration at saturation of about 2000 ppm.
Figure 10 shows the shear strain amplitude dependence of storage and loss moduli of
␾ = 0.45 kaolinite suspensions 共i兲 without polymer 共ii兲 with 2000 ppm of HPG, corresponding to the estimated polymer concentration at saturations 共iii兲 with 6250 ppm of
HPG, i.e., with polymer in excess. From a qualitative point of view, the shear strain
amplitude dependence of viscoelastic moduli of suspensions in the presence of polymer is
FIG. 10. Shear strain amplitude dependence of storage modulus 共open symbol兲 and loss modulus 共full symbol兲
of kaolinite suspensions at ␾ = 0.45 without polymer 共〫 , ⽧兲, with 2000 ppm 共䊐 , 䊏兲 and 6250 ppm 共䊊 , 쎲 兲 of
HPG.
1266
FIG. 11. Storage modulus in the linear regime, loss modulus in the linear regime and critical shear strain
amplitude ␥c of kaolinite suspensions at ␾ = 0.45 as a function of polymer concentration, ␻ = 1 Hz.
similar to that observed without polymer. It is worth mentioning that shear strain amplitude ␥max increases drastically with increasing polymer concentration, especially above
2000 ppm. Linear viscoelastic parameters G0⬘, G0⬙ and ␥c are plotted as a function of
polymer concentration in Fig. 11. As polymer concentration increases up to 2000 ppm,
both G0⬘ and G0⬙ moduli increase sharply and the critical shear strain amplitude remains
nearly constant. Above 2000 ppm, both G0⬘ and G0⬙ moduli decrease and ␥c increases
drastically. The dependence of rheological parameters with polymer concentration in the
linear viscoelastic domain is shown in Fig. 12. The cohesive energy density curve 关Fig.
12共a兲兴 displays an intermediate concentration regime, from 2000 to 4000 ppm, where it is
constant, separating two regimes at low and high concentrations, characterized by a
substantial increase of cohesion energy density with increasing polymer concentration.
Therefore, the concentration of ⬃2000 ppm, corresponding to the polymer concentration
at saturation, appears as a transition concentration in the rheological behavior of concentrated suspensions. The description in three concentration regimes is also relevant regarding the reduced intensity of the G⬙ peak, presented in Fig. 12共b兲. In the lowest concentration regime, that is up to 2000 ppm, the intensity of the strain hardening in G⬙
increases sharply upon increasing polymer concentration. The intermediate concentration
regime is characterized by a significant reduction of the rate of increase of the Gmax
⬙ / G⬙
ratio, up to 4000 ppm. Above 4000 ppm, Gmax
⬙ / G⬙ decreases extensively, and finally
tends to a nearly constant value at high polymer concentrations.
From a molecular point of view, adsorbed polymer chains are known to form a polymer layer with a thickness close to the radius of gyration, Rg, of free chains 关Semenov et
al. 共1996, 1997兲兴. Assuming that kaolinite particles are uniformly dispersed, the average
inter-particle distance is about 40 nm at ␾ = 0.45, which is much smaller than 2Rg
⬃ 180 nm, meaning that adsorbed polymer chains are most likely confined within this
narrow inter-particle region. As a consequence, in the low concentration regime, the
contribution of polymer chains to the enhancement of both cohesive energy density and
strain hardening in G⬙ can be attributed to both polymer bridges between clay particles
and steric interactions of confined adsorbed polymer. However, at high polymer concentrations, above 4000 ppm, the enhancement of cohesive energy density is mostly due to
steric interactions induced by free polymer chains, which make particle bridges less
1267
FIG. 12. 共a兲 Cohesive energy density Ec and 共b兲 reduced intensity of the G⬙ peak of kaolinite suspensions at
␾ = 0.45 versus polymer concentration, ␻ = 1 Hz.
likely. The rate of increase of Ec, which is more important in the lower concentration
regime than that in the higher one, suggests that the contribution of bridging polymer
chains to cohesive energy density is more important than that induced by steric interactions between polymer chains, which makes sense. Concurrently, free chains induce
lubrication effects between kaolinite particles covered by adsorbed polymer chains,
which may be responsible for the sharp decay of the strain hardening in G⬙ above
4000 ppm, and could also explain the high value of the shear strain amplitude, ␥max,
corresponding to the maximum of the strain hardening in G⬙. In the intermediate concentration regime, steric interactions between adsorbed polymer layers overcome contributions to cohesive energy density from particle bridging, finally leading to a leveling off
of Ec.
IV. CONCLUDING REMARKS
Linear and nonlinear viscoelastic behaviors of concentrated suspensions of kaolinite
plate-like particles have been investigated as a function of volume fraction, ionic
strength, and polymer concentration. These three adjustable parameters have been used to
tune the influence of inter-particle interactions: excluded volume interactions between
clay particles, long range electrostatic repulsive interactions and steric interactions mediated by adsorbed polymer chains.
1268
As repulsive interactions increase, by either increasing volume fraction or increasing
polymer concentration or decreasing ionic strength, the cohesive energy density of concentrated suspensions increases.
The nonlinear viscoelastic behavior has been thoroughly investigated. In particular, the
G⬙ vs. strain amplitude curve exhibits a hump, characterized by a relative intensity which
can be tuned by repulsive interactions. This strain hardening in G⬙ is mostly enhanced by
increasing repulsive interactions, except at high polymer concentrations, due to lubrication effects mediated by free polymer chains.
The whole set of rheological data suggests that the strain hardening in G⬙ originates
from extra electro-viscous dissipation, due to the coupled increase of excluded volume
and electrostatic interactions, which are both modulated by the fluctuation of the interparticle distance during oscillatory shear measurement.
Appendix
Under oscillatory shear, any vector du共dx0 , dy 0兲 is transformed into du⬘共dx0⬘ , dy 0⬘兲:
共du⬘兲 = 共F0兲共du兲 or
冉冊冉
dx0⬘
dy 0⬘
=
dx0 + ␥0 · dy 0
dy 0
冊
.
In the coordinate system of the laboratory, 共x0 , y 0兲, the gradient tensor of the transformation is given by
1 ␥0
共F0兲 =
0 1
.
In the coordinate system of the particle 共x , y兲, the gradient tensor of the transformation
is given by 共F兲 = 共R共−1
␣,z0兲兲共F0兲共R共␣,z0兲兲.
With
cos ␣ sin ␣
共R共␣,z0兲兲 =
,
− sin ␣ cos ␣
consequently:
冉冊
冉
共F兲 =
冢
1−
␥0
sin 2␣
2
␥0 sin ␣
2
冊
␥0 cos2 ␣
␥0
1+
sin 2␣
2
冣
= 共1d兲 + 共ⵜ␰兲,
with 共Id兲 the identity matrix and 共ⵜ␰兲, the displacement gradient tensor
共ⵜ␰兲 =
冢
−
␥0
sin 2␣ ␥0 cos2␣
2
␥0
␥0 sin2␣
sin 2␣
2
冣
.
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Rheological Characterization of Starch
Derivatives/Polycaprolactone Blends
Processed by Reactive Extrusion
Frédéric Bossard,1 Isabelle Pillin,2 Thierry Aubry,3 Yves Grohens2
1
Laboratoire de Rhéologie, Grenoble Institut National Polytechnique, Université Joseph Fourier Grenoble I,
UMR 5520, BP 53, 38041 Grenoble Cedex 9, France
2
Laboratoire Polymères, Propriétés aux Interfaces et Composites, Université de Bretagne Sud,
Rue de Saint Maudé, BP92116, 56321 LORIENT CEDEX, France
3
Laboratoire de Rhéologie, Université de Bretagne Occidentale, 6 avenue Le Gorgeu,
CS 93837, 29238 BREST CEDEX 3, France
Linear viscoelastic, steady shear behaviors, and morphologies of starch formate/poly(e-caprolactone) (PCL)
blends, compatibilized by oligomers and obtained by
reactive extrusion, have been investigated as a function of formic acid (FA)/starch ratio, nature, and molecular weight of the oligomer. The rheological properties
of these blends have been compared with those of a
commercial starch-based thermoplastic, namely MaterBi1 ZF03UA. In presence of FA, starch is destructured
to starch formate and oligomers are used as plasticizers. The linear viscoelastic response of blends is quite
similar to that of nanocomposite materials: the low frequency behavior is attributed to a percolated network
of destructured starch particles, and the high frequency behavior is that of the polymer matrix. The viscosity curve presents a profile characterized by two
plateau regions, at low and high shear rates. The plateau region at low shear rates corresponds to the viscous response of the blend while that observed at high
shear rates can be attributed to the PCL matrix. The
compatibilization is enhanced in the presence of starch
formate and increases with increasing the oligomer
molecular weight. The use of PCL oligomer was shown
to improve this compatibilization effect. POLYM. ENG.
SCI., 48:1862–1870, 2008. ª 2008 Society of Plastics Engineers
INTRODUCTION
Replacement solutions of petroleum-derived plastics by
biopolymers have focused increasing attention in the past
Correspondence to: Frédéric Bossard; e-mail: [email protected]
DOI 10.1002/pen.21160
Published online in Wiley InterScience (www.interscience.wiley.com).
C 2008 Society of Plastics Engineers
V
few years, even if they raise new economical problems
such as an increasing price of food breeds. Amongst such
biomaterials, starch-based polymer blends have attracted
much attention, mainly because starch is an inexpensive
and abundant biopolymer derived from renewable resource. Studies focus mainly on the compounding processes [1], the influence of plasticizers [2], and their relation with rheological [3–6] and thermomechanical properties [7–9]. Native starch is constituted of linear amylose
and branched amylopectin molecules with micrometric
granular shape. The complex structure of granular starch
results from the succession of amorphous and crystalline
growth rings, which is related to the molecular organization of amylopectin. The branched glycan chains of the
amylopectin form an amorphous phase at the root of each
unit cluster while they are organized in double helical
structures far from this root, responsible for the crystalline
sections [10–12]. Native starch can be directly blended
with polymer matrix but it is incompatible with most
thermoplastics and remains generally in a micrometric
granular form, leading to heterogeneous materials known
to exhibit poor compressibility, flexibility, and elasticity
properties [13–15]. The heterogeneous character of blends
can be reduced by destructuring native starch using plasticizers such as water, glycerol, or sorbitol [16]. This modification, favored by heating, consists in a limited disbranching of amylopectin but also a destructuration of the
crystalline network that leads to the swelling and gelatinization of starch.
A large number of studies deals with association of
native starch, or modified starch, with a biopolymer such
as poly(e-caprolactone) (PCL), poly(lactic acid) (PLA),
poly(hydroxyalcanoate) (PHA), poly(butylene adipate
POLYMER ENGINEERING AND SCIENCE—-2008
terephthalate) (PBAT) [17–23]. An amount of starch,
ranging from 40 to 60% is generally added to polymeric
matrix, leading to the formation of starch composites
called plastified starch materials. Novamont, S.p.A.,
Novara (Italy) commercializes various starch-based composites under the trademark Mater-Bi1. Their mechanical
properties are similar to conventional synthetic thermoplastics, such as polypropylene [24], and they are generally resistant to oils and alcohols; however, their high production cost is a major drawback, which limits their use
in industrial applications.
Therefore, research works were carried out to provide
new blends of starch/thermoplastic materials exhibiting
high rheological properties. Following this objective, we
have recently studied the use of synthetic biopolymers,
such as PCL, and various polyester oligomers as compatibilizers and plasticizing agents of both starch and PCL
[25, 26]. Indeed, the plasticizing effect of oligomers has
been pointed out by the significant swelling of starch
granules combined with a decrease of the Tg of PCL on
adding low molecular weight oligomers. In this previous
work, a new chemical process of destructuration of starch
to starch formate with formic acid (FA) was proposed.
The O-formylation reaction of starch is a reversible reaction leading to the rapid formation of monoformic ester,
and the substitution rate, controlled by the FA/starch ratio,
can reach 1.2 [27]. Starch granules can be fully destructured by concentrated FA solution in a batch reactor. FA
destructuration of starch means destructuration of native
starch granules but not depolymerization, which occurs
during starch hydrolysis. Blends were obtained by mixing
starch formate to PCL matrix, and the use of starch formate in the formulation of starch/PCL thermoplastic
blends was shown to be of interest for applications. Moreover, in these previous studies, we have shown that PCL
oligomer was the most efficient compatibilizer for this
specific blend, because PCL oligomer has a good miscibility with PCL matrix and exhibits strong interactions
with modified starch.
However, this two-step process, i.e., (i) degradation of
native starch into starch formate and (ii) extrusion of
starch formate/PCL blends, is not a suitable process for
industrial applications and needs to be adapted to standard
processing equipments. That is the reason why, in the
present article, the starch formylation was included in a
one step reactive extrusion process. The ratio FA/starch
and the nature of the oligomer have been chosen as parameters for the formulation. Linear and nonlinear viscoelastic investigation techniques, combined with scanning
electron microscopy, have been used to study the relation
between rheological properties and morphology of blends.
Attention has been paid to the influence of FA/starch ratio
and the nature of the oligomer on composite morphologies and their rheological properties. Results obtained
with this new process are compared to those obtained
with a Mater-Bi starch-based materials named ZF03UA
and commercialized by Novamont S.p.A.
DOI 10.1002/pen
MATERIALS AND METHODS
Materials
A commercial starch-based composite, named MaterBi ZF03UA and supplied by Novamont S.p.A (Italia), has
been used as a reference material in this study. ZF03UA
grade contains about 40 wt% of destructured wheat starch,
blended within a PCL matrix. Characterized by a melting
temperature of 648C [28], a glass transition temperature
of about 2608C, a strength at break of 31 MPa and an
elongation at break of 886% [29], Mater-Bi Z grades are
especially suitable for film blowing process.
In this study, starch-based blends are composed on
wheat starch (I59-113H10), purchased from ROQUETTE
(France). The water content in native starch, measured by
TGA, is about 13 wt%. FA solution at 99 wt% was purchased from Sigma Aldrich and PCL, referenced as
CAPA 6800, was provided by SOLVAY. Its molecular
weight is 80,000 g mol21, corresponding to a density of
1.11 g cm23. Three polyester oligomers, provided by
DUREZ, were used as plasticizers: two 1,6-hexane-diol
adipate and phthalates, referenced as 105-42 and 105-15
with different molecular weights, and a PCL bearing
hydroxyl functions, referenced as 1063-35. The oligomers
were selected because of their molecular structure and
low molecular weight. The low molecular weight was
chosen to decrease the entropic negative contribution of
the mixing free energy. The polyester structure was chosen as close as possible to that of the PCL matrix, to get
good compatibility. Main characteristics of the three
oligomers are presented in Table 1. Three blends, named
B1, B2, and B3, were studied; all composed of 40 wt%
starch with a various FA ratio, 30 wt% PCL and 30 wt%
oligomer. Formulations differ from the nature of the
oligomer and the FA ratio: B1 and B2 contain the
oligomer 105-42 and 105-15, respectively, while B3 is
composed of PCL 1063-35. In the text, the blend name is
followed by the FA/starch ratio, calculated with respect to
the starch amount in the extruder. So that, the blend
assigned as B1-30 is composed of phthalate 105-42 and
contains 40% of starch and FA with FA/starch ratio of
30 wt%. Blend compositions are summarized in Table 2.
Methods
Sample Processing. The blends were extruded with a
BC21 CLEXTRAL (Firminy, France) corotating twinscrew industrial scale extruder controlled by a Lab-station. The screw geometrical features were the following:
25-mm diameter, 900-mm length, 21-mm axial length.
The screws configuration was: 0–100 mm: T2F (trapezoidal double thread); 100–625 mm: C1F (conjugated single
thread); 625–850 mm: C2F (conjugated double thread);
850–862.5 mm: C1FC (conjugated single thread, with
direction of threading contrary to the III configuration);
POLYMER ENGINEERING AND SCIENCE—-2008 1863
TABLE 1. Main characteristics of DUREZ oligomers.
Oligomer
Nature
Functionalization
Mw (g mol21)
105-42
105-15
1063-35
1,6-hexane-diol adipate and phthalate
1,6-hexane-diol adipate and phthalate
Poly(e-caprolactone)
Hydroxyl
Hydroxyl
Hydroxyl
2700
7400
2000
862.5–875 mm: C1F (conjugated single thread); 875–900
mm: C2F (conjugated double thread). The extruder is
composed of nine thermo-controlled elements with the
following profile of temperature: 20-20-20-22-40-60-7090 and 1008C, and the screw rotation speed was fixed at
350 rpm. Starch-based thermoplastics were obtained by
blending first starch granules with the PCL matrix in the
hopper. Then, FA was injected in the third and fourth
thermo-controlled elements, that is, at 20 and 228C,
respectively. The oligomers were added between the fifth
and the sixth element temperature that is at 40 and 608C,
respectively. The mechanical work during extrusion is
given by the specific mechanical energy (SME), calculated using Eq. 1 [30]:
SME ðkJ=kgÞ ¼
2pNT
1000m
(1)
where N is the screw rotation speed in revolution per second, T is the torque in Nm, and m is the mass flow rate
in kg/s. The extruded blends were quenched in cold
water. After being dried 12 h in deep vacuum to prevent
water uptake, blends were compression-molded between
2-mm-thick sheets, using a press under 20 MPa at 1408C.
Rheological Measurements. The linear and nonlinear
rheological measurements were carried out using a stress
controlled Malvern Instruments Gemini rheometer,
equipped with a parallel-plate geometry (diameter ¼ 25
mm, gap ¼ 2 mm). Measurements were performed under
continuous purge of dry nitrogen to avoid thermo-degradation and moisture adsorption and at a temperature of
958C, that is, 408C above the melting temperature of
blends B1, B2, and B3 and 318C above the melting tem-
perature of ZF03UA. For oscillatory shear measurements,
a strain amplitude of 0.1% was applied, which is well
below the linear viscoelastic limit for all samples. Creep
measurements have been used to determine the flow
curves at low shear rates. The linear time-dependence of
strain has been verified over more than 1500 s, so that
steady state was achieved for each creep measurement.
The degradation state of native starch is monitored
through reduced viscosity measurements of starch formate
using a capillary Schott Gerate Ubbelohde tube driven by
a SCHOTT GERATE AVS 310 controller. A starch solution at a concentration of 0.5 wt% is obtained by dispersing starch formate in dimethylformamide (DMF) with
5 wt% LiCl. The reduced viscosity of starch formate is
deduced from the measurement of the flow time of the
solution, t, and that of the pure solvent, t0, as follows:
Zred ¼
Z Z0 t t0
¼
CZ0
Ct0
(2)
where C is the polymer concentration (g ml–1).
Scanning Electron Microscopy. Morphological observations were performed using a JEOL JSM-6031F scanning electron microscope. Cryofractured sheets of 2-mmthick were vacuum-metalized before observation.
RESULTS AND DISCUSSION
Mater-Bi ZF03UA
Figure 1 shows a scanning electron microscopy image
of cryogenic fractures of ZF03UA. The commercial
starch-based composite exhibits a continuous phase mor-
TABLE 2. Composition of the blends as a function of starch, oligomers, PCL weight percentages and FA/starch ratio.
Reference
B1
B2
B3
0
15
30
60
0
15
30
60
0
15
60
FA/starch ratio%
Oligomer
Starch %
PCL %
Oligomer %
0
15
30
60
0
15
30
60
0
15
60
105-42
40
30
30
1864 POLYMER ENGINEERING AND SCIENCE—-2008
105-15
1063-35
DOI 10.1002/pen
FIG. 1. SEM micrograph of cryo-fractured surface of ZF03UA MaterBi1: continuous phase of PCL with dispersion of starch.
phology of PCL without visible starch nodules at a micrometric scale, which suggests a destructuration of starch at
a sub-microscopic level, due to the gelatinization and
complexation process. Figure 2 shows the storage and loss
moduli versus frequency for ZF03UA samples. Both moduli are nearly constant at the lowest frequencies investigated, while they increase significantly with increasing
frequency above x ¼ 1 rad s21. In the case of ZF03UA,
amylopectin nanoparticles are chemically extracted from
starch and then recombined with an amylose/surfactant
complex to improve the compatibilization with the polyester matrix. The compatibilized blend of fully destructured starch particles and PCL confers to ZF03UA high
viscoelastic moduli. A solid-like response at low frequencies is generally observed for starch-based thermoplastic
and has been attributed to a reduction of the mobility of
the amorphous phase induced by the presence of remaining semi-crystalline domains arising from the formation
of amylose-lipid complexes [5]. However, the formation
of a network of connected destructured starch could also
explain such solid-like behavior. The viscoelastic behavior
FIG. 3. Steady shear viscosity (open symbol) versus shear rate and
complex viscosity (full symbol) versus frequency of Mater-Bi1 ZF03UA
at 958C. The dash/dot lines represent fits of the two contributions g1 and
g2 using Cross model and the solid line represents the sum of the fits.
at high frequencies can be reasonably attributed to the
viscoelastic response of the PCL matrix.
Figure 3 presents the complex viscosity of ZF03UA as
a function of frequency (full symbols), extended towards
low shear rates by creep measurements (open symbols).
Surprisingly, it has to be noticed that the empirical CoxMerz rule, which states that the complex viscosity as a
function of frequency is equal to the apparent viscosity as
a function of shear rate, is satisfied. Indeed, the Cox-Merx
rule has shown to describe properly flow properties of
melting polymers and polymer solutions but it is not generally suitable for charged polymer. The viscous profile of
ZF03UA can be considered as the superposition of two
_ and g2 ðcÞ,
_ associated with a very high
contributions, g1 ðcÞ
plateau viscosity, g01, at low shear rates, and to a moderate plateau viscosity, g02, at high shear rates. The existence of a very high zero-shear viscosity, followed by
drastic shear-thinning, characterized by a power law shear
rate dependence of the apparent viscosity with an exponent close to 21, is indicative of an apparent yield stress
behavior at intermediate shear rates. The plateau viscosities g01 and g02 have been obtained by fitting the experimental data, at low and high shear rates respectively,
using the Cross model:
ZðġÞ ¼
FIG. 2. Storage modulus G0 and loss modulus G00 of Mater-Bi1
ZF03UA at 958C.
DOI 10.1002/pen
Z0 Z1
þ Z1
1 þ ðKġÞn
(3)
in which Z‘, K and n are taken as adjustable parameters
in this work. Let us stress that the sum of those two fits
describes properly the whole profile of viscosity. The
value of Z02 ¼ 5 104 Pa.s is consistent with the zeroshear viscosity of the PCL matrix having a molecular
weight of about 1.2 105 [31]. This result points out that
the flow curve at high shear rates is dominated by the viscous response of the PCL matrix, in agreement with linear
viscoelastic data. The very high plateau viscosity at low
FIG. 4. SEM micrographs of cryo-fractured surface of B1 blends (a) without formic acid (B1-0); (b) with a
FA/starch ratio of 15% (B1-15); (c) 30% (B1-30), and (d) 60% (B1-60).
shear rates, Z01 ¼ 8 107 Pa.s, associated with the existence of an apparent yield stress, could be attributed to
the viscous contribution of the percolated network of destructured starch, as suggested by the previously-discussed
linear viscoelastic response. Indeed, the presence of
remaining semi-crystalline domains cannot explain the
yield stress response of the blend.
Starch Formate-Based Thermoplastics
Influence of the FA/Starch Ratio. Figure 4a–d show
scanning electron micrographs of blends B1-0, B1-15,
B1-30, and B1-60, respectively. Contrary to the structure
of the Mater-Bi ZF03UA, B1 blends contain micrometric
starch particles. The presence of micrometric starch globules in B1 blends is due to softer processing conditions in
terms of temperature and mechanical treatment, compare
to those used to obtain ZF03UA. Indeed, with an SME of
about 300 kJ/kg, the mechanical input is not strong
enough to produce a destructuration of starch at the molecular scale and lead to granule fragmentation [32]. The
processing conditions to obtain B1 differ also from that
used in a batch reactor by the smaller amount of FA
added and the fact that blends have been obtained by
extrusion process. Consequently, some starch granules are
expected to be unaffected by the FA treatment. However,
it is difficult to assess the ratio of non-modified starch
granules. Moreover, FA has an effect on PCL degradation
too. Indeed, a decrease of PCL molecular weight due to
ester groups’ hydrolysis has been observed above 1008C
under acidic conditions [33]. Without FA, starch particles
appear to be well dispersed within the PCL matrix and
not destructured by the thermomechanical treatment
imposed during extrusion. In the presence of an increasing
amount of FA, at least up to a FA ratio of 30%, starch
nodules do not seem to be modified in shape and starch
nodules of B1-60 blend appear to be connected. Reduced
viscosity measurements of starch have been carried out to
monitor the influence of FA on the destructuration state
of starch. For this purpose, starch granules were extruded
in the same conditions than the blends without and with
15 and 30% of FA/starch ratio and dispersed in DMF
with 5% LiCl. Table 3 shows that the reduced viscosity
of such starch dispersions is nearly unmodified when
increasing FA/starch ratio, suggesting that FA does not
modify significantly the molecular weight of amylose and
amylopectin molecules. Still it has to be noticed that the
gradual increase of FA/starch ratio induces small cracks
in the matrix, which could be the mark of a chemical
attack of PCL polymer.
TABLE 3. Reduced viscosity of starch dispersed in 5% LiCl DMF at
various FA/starch ratios.
% Starch
100
85
70
% FA/starch
Reduced viscosity (ml/g)
0
15
30
220
218
210
DOI 10.1002/pen
FIG. 6. Flow curves of B1 blends without formic acid (B1-0) and with
FA/starch ratios of 15% (B1-15), 30% (B1-30), and 60% (B1-60) at
958C. Open symbols represent creep measurements while full symbols
represent dynamic measurements.
FIG. 5. (a) Storage modulus G0 ; (b) loss modulus G00 of B1 blends
without formic acid (B1-0) and with FA/starch ratios of 15% (B1-15),
30% (B1-30), and 60% (B1-60) and the matrix (solid line) as a function
of frequency at 958C.
Figure 5a and b show the frequency dependence of
storage and loss moduli of blends B1, at various FA/
starch ratios at 958C. From a qualitative point of view,
the G0 behavior of B1 blends is significantly influenced
by FA concentration, contrary to G00 , whose qualitative
behavior seems to be nearly insensitive to FA treatment.
More precisely, for all FA/starch ratios considered in this
work, G0 exhibits a trend to a plateau value at low frequencies and tends to the G0 of the matrix at high frequencies, as observed for Mater-Bi ZF03UA samples.
Besides, for FA/starch ratios above 15%, the low frequency G0 plateau is all the less marked as FA/starch ratio
is higher, suggesting that optimum structuration is
achieved at FA/starch ratio 15%. Nevertheless, this
structuration effect is certainly far less marked for B1
samples than for Mater-Bi ZF03UA samples, as evidenced
by the fact that G0 and G00 values have much lower values
than for Mater-Bi ZF03UA and no low frequency G00 plateau could ever be observed for any B1 samples. The high
level of linear viscoelastic moduli of Mater-Bi ZF03UA
compared to that of B1 samples is probably due to the total
melting and destructuration of starch in the Mater-Bi
ZF03UA samples compared to the crude dispersion of starch
microparticles in the B1 samples, however an enhancement
of starch/PCL compatibilization cannot be excluded.
Figure 6 shows the viscosity as a function of shear
rates as well as the complex viscosity as a function of
DOI 10.1002/pen
frequency for B1-0, B1-15, B1-30, and B1-60. All blends
exhibit a flow curve similar to that obtained for Mater-Bi
ZF03UA samples, still with a lower zero-shear viscosity
g01 and a more marked second viscosity plateau g02. Concerning B1 blends, two distinct effects of FA are pointed
out in Fig. 6: the zero-shear viscosity of the blend, g01,
reaches a maximum for a FA/starch ratio 15% whereas
the matrix viscosity, g02, decreases gradually with increasing FA/starch ratio.
The above-described dependence of viscoelastic moduli
and plateau viscosities as a function of FA/starch ratio
can be interpreted as a combined effect of the FA on
starch and on the matrix. First, FA induces a partial
destructuration of starch particles, most probably at the
starch granule surface, with the formation of starch formate. An increasing FA/starch ratio favors compatibilization of starch with PCL through an increase of the interactions between hydroxyl and ester functions of the
oligomer and hydroxyl and formate functions of modified
starch, leading to a reinforcement of the blend rheological
properties, but a weakening of the matrix ones. A FA/
starch ratio of about 15% appears to be the optimum
ratio, at least regarding the rheological response of this
blend.
Influence of Oligomer Molecular Weight. In this section, we have investigated B2 blends that differ from B1
grades in the molecular weight of the oligomer, passing
from 2700 g/mol for B1 to 7400 g/mol for B2. Microstructures of B2 blends B2-0, B2-30, and B2-60 are
shown in Fig. 7a–c, respectively. All blends exhibit a
microstructure quite similar to that of B1 blends, suggesting a similar mechanical fragmentation of starch granules.
However, upon increasing the amount of FA, the interface
between starch nodules and PCL matrix is less regular
and starch seems to be more connected than for B1
FIG. 7. SEM micrographs of cryo-fractured surface of (a) B2-0, (b) B2-30, (c) B2-60, and (d) B3-15.
blends, leading to a nearly continuous phase structure for
B2-30 and B2-60 blends.
Figure 8a and b show storage and loss moduli as a
function of frequency for B2 blends at various FA/starch
ratios at 958C. With increasing FA/starch ratio, the tendency to form a G0 plateau at low frequencies is more pronounced than for B1 samples and is maximum for a FA/
starch ratio of about 30%, instead of 15% for B1 samples.
This difference is most likely due to a more connected
structure for B2 blends since starch destructuration, and
consequently cristallinity of blends, is expected to be
unchanged when oligomer molecular weight is changed.
These differences between B1 and B2 samples are confirmed in Figure 9, showing the viscous behavior of B2
blends at various FA/starch ratios. Indeed both plateau
viscosities, that is blend viscosity g01 and matrix viscosity
g02, reach a maximum for a FA/starch ratio of about
30%, which again appears as the optimum ratio for B2
blend. The high value of zero shear viscosity observed for
B2 blends could be attributed to the reduction of the solubility of oligomers in starch formate when increasing the
oligomer molecular weight [34]. Indeed, the longer
oligomer molecules are expected to be localized:
at the formate/PCL interface, favoring interactions
between starch nodules, and therefore enhancing
connectivity between nodules,
within the PCL matrix, thus contributing to the
enhancement of viscosity of the matrix.
FIG. 8. (a) Storage modulus G0 ; (b) loss modulus G00 of the matrix
(solid line), and of B2 blends without formic acid (B2-0), with a FA/
starch ratio of 15% (B2-15), 30% (B2-30), and 60% (B2-60), as a function of frequency at 958C.
DOI 10.1002/pen
FIG. 9. Flow curves of B2 blends without formic acid (B2-0) and with
a FA/starch ratio of 15% (B2-15), 30% (B2-30), and 60% (B2-60) at
958C. Open symbols represent creep measurements while full symbols
represent dynamic measurements.
Influence of the Nature of the Oligomer. The key role
played by the oligomer, acting as a compatibilizer in a
batch reactor, has been shown in a previous paper [26]. In
this reference, the comparison between relaxation times of
starch formate/PCL and starch/PCL blends has been performed using solid-state NMR investigation techniques.
The modification by FA has been shown to yield a
decrease of the difference of the relaxation times between
starch and PCL characteristic chemical groups, proving
that the formylation of starch improved the miscibility of
starch with PCL. In order to increase further the compatibilization, the 1,6-hexane-diol adipate and phthalate
oligomer has been substituted in the B3 blends by PCL
oligomers that have a better affinity with starch formate
than with starch. Such enhanced compatibility is due to
intermolecular hydrogen bonds between PCL oligomer
and starch formate [26]. Moreover, obviously PCL
oligomer has a natural compatibility with the PCL matrix.
Figure 10a and b show the frequency dependence of G0
and G00 moduli of B3-0, B3-15, and B3-60. First, contrary
to B1 and B2 blends containing the 1,6-hexane-diol adipate and phthalate oligomer, the tendency to form a G0
plateau at low frequencies for B3 blends is pronounced
even without FA. This difference between B3 and both
B1 and B2 blends is only ascribable to a better compatibilization of starch formate with the PCL matrix, induced
by the improved affinity of the PCL oligomer with both
elements. This interpretation is confirmed by SEM micrographs of B1-15 in Fig. 7d, in which no starch particles
are perceptible, suggesting a pronounced compatibilization of starch with the matrix during the reactive extrusion. At low frequencies, linear viscoelastic moduli
increase with increasing FA/starch ratio up to 15% and
then strongly decrease with further addition of FA. At this
optimum FA/starch ratio of 15%, G0 and G00 moduli reach
a high value of about 104 Pa at low frequencies, that is,
DOI 10.1002/pen
FIG. 10. (a) Storage modulus G0 ; (b) loss modulus G00 of the matrix
(solid line), and of B3 blends without formic acid (B3-0) and with a FA/
starch ratio of 15% (B3-15) and 60% (B3-60), as a function of frequency
at 958C.
much higher than those measured for B1 and B2 blends,
and of the order of magnitude of that for Mater-Bi
ZF03UA samples. This solid-like response at low frequency can be explained by the expected higher connected
structure induced by the PCL oligomer, as depicted by
SEM micrographs. This result is in accordance with
flow curves of B3 blends, presented in Fig. 11. Indeed the
FIG. 11. Flow curves of B3 blends without formic acid (B3-0) and
with a FA/starch ratio of 15% (B3-15) and 60% (B3-60) at 958C Open
symbols represent creep measurements while full symbols represent
dynamic measurements.
zero-shear viscosity of B3-15 far exceeds that of B3-60,
and is not so far from that for Mater-Bi ZF03UA samples.
CONCLUSION
In the present article, linear viscoelastic and steady
shear measurements performed on new biopolymer starchbased PCL blends have been presented and compared to
the rheological properties of a commercial starch-based
polymer, namely Mater-Bi ZF03UA.
We have first demonstrated the ability to provide bioplastic starch-based PCL blends, exhibiting high zeroshear viscosity and linear viscoelastic moduli in the melt.
Using FA, starch formate was formed through an O-formylation reaction, which was included in a one step
reactive extrusion process. The whole set of rheological
measurements has underlined the key role of the
oligomer, acting as a compatibilizing agent between the
starch formate and the PCL matrix. The addition of such
oligomers leads to an optimum of the blend rheological
properties for a FA/starch ratio ranging from 15 to 30%,
depending on the oligomer molecular weight. The use of
high molecular weight oligomers having a strong affinity
with both starch formate and the matrix significantly
improves the blend rheological properties.
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DOI 10.1002/pen
Rheol Acta (2010) 49:529–540
DOI 10.1007/s00397-009-0402-8
Author's personal copy
ORIGINAL CONTRIBUTION
Influence of dispersion procedure on rheological properties
of aqueous solutions of high molecular weight PEO
Frédéric Bossard · Nadia El Kissi ·
Alessandra D’Aprea · Fannie Alloin ·
Jean-Yves Sanchez · Alain Dufresne
Received: 30 April 2009 / Accepted: 18 November 2009 / Published online: 11 December 2009
© Springer-Verlag 2009
Abstract The linear and nonlinear viscoelastic behaviors of poly(ethylene oxide) (PEO) in aqueous media
have been investigated as a function of concentration
and molecular weight. A particular interest has been
paid to study the effect of turbulent flow under stirring, inducing both shear and elongational stresses,
on the rheological behavior of the polymer solutions.
The comparison of intrinsic viscosity and viscoelastic
properties between shaken and stirred PEO solutions is
discussed at the molecular scale in terms of chain scis-
sion and aggregation. Results point out that the effect
of the mechanical history on the rheological response
of PEO solutions depends also on the concentration
regime and molecular weight. Indeed, the influence of
the dispersion procedure vanishes by decreasing both
the concentration and the molecular weight.
Keywords Poly(ethylene oxide) · Chain scission ·
Aggregation · Molecular weight · Linear viscoelasticity
Introduction
Paper presented at the de Gennes Discussion Conference
held February 2–5, 2009 in Chamonix, France.
F. Bossard (B) · N. El Kissi · A. D’Aprea
Laboratoire de Rhéologie,
Grenoble Institut National Polytechnique,
Université Joseph Fourier Grenoble I,
UMR 5520, BP 53, 38041 Grenoble, France
e-mail: [email protected]
F. Alloin · J.-Y. Sanchez
LEPMI, Grenoble Institut National Polytechnique,
Université Joseph Fourier Grenoble I,
CNRS, UMR 5631, 1130 rue de la piscine,
BP 75, 38402 Saint Martin d’Hères, France
A. Dufresne
Ecole Internationale du Papier,
de la communication imprimée et des Biomatériaux,
PAGORA- Grenoble Institut National Polytechnique,
BP 65, 38041 Grenoble, France
Present Address:
A. Dufresne
Departamento de Engenharia Metalurgica e de Materiais,
Universidade Fédéral do Rio de Janeiro (UFRJ),
Coppe, Rio de Janeiro, Brazil
Poly(ethylene oxide) (PEO) has been extensively studied due to its unique behavior in aqueous media and
its important industrial application. Its biocompatibility
and protein adsorption inhibitor capacity make PEO
a good candidate for the development of a new drug
delivery medicine (Lee et al. 1995; Allen et al. 1999),
and it can be used to substitute some biopolymers
as implant for tissue replacement or augmentation
(Villain et al. 1996). The specific chemical structure
of PEO, HO − [(CH2 )n − O]x − H with n = 2, confers
to this polymer very unusual interactions with water.
Indeed, while poly(methylene oxide) with n = 1 and
poly(butylenes oxide) with n = 3 are both hydrophobic
and insoluble in water, PEO is known as the simplest hydrosoluble hydrocarbon polymer, regarding its
chemical structure. Its solubility in water originates
from the competition between PEO–water and water–
water hydrogen bonding (Dormidontova 2002, 2004),
delicately balanced by hydrophobic interactions induced by the ethylene components. The rupture of
hydrogen bonds with increasing temperature is responsible for the decrease of its solubility upon heating,
530
a lower critical solution temperature behavior. Near
room temperature, the water solubility of PEO is found
to depend also on the polymer concentration. Indeed,
water is a good solvent at low concentration and high
temperature while it becomes a bad solvent at intermediate concentration, close to the critical concentration
(Daoust and St Cyr 1984).
PEO is used for applications requiring high cation
solvation and good electrochemical stability such as
solid electrolyte used in lithium polymer cells (Gray
and Armand 2000; Wright 1998), but surprisingly, it
appears to be fragile and very sensitive to thermal
(Vijayalakshmi et al. 2005), photochemical (Morlat and
Gardette 2003; Hassouna et al. 2007), and ultrasound
(Pritchard et al. 1966; Kanwal and Pethrick 2004)
degradations in the bulk and in solution. For polymer solutions, thermal and photochemical degradations have been observed respectively for temperatures
higher than 50◦ C and for samples exposed to irradiation
corresponding to natural outdoor aging (λ > 300 nm).
Both degradation processes induce the formation of
formate and ester groups. The release of formic acid
ions (HCOO− ) is responsible for a drastic decrease
of the pH, leading to a random chain scission. In the
case of ultrasound degradation, chain scission is due
to intense shear stresses arising from the collapse of
transient cavitation bubbles. Contrary to thermal and
photochemical degradations, the resulting rupture of
covalent bonds occurs preferentially in the middle of
polymer chains (Madras and McCoy 2001) up to a
lower limit of molecular weight of about 2 × 104 g/mol
below which the polymer will not undergo scission.
It has to be noticed that sonication produces heat,
leading to some local increases of temperature and consequently associated thermodegradation, even if the
temperature of the sample is externally controlled.
Consequently, PEO in water requires controlled and
precise conditions when handling, which make it a
delicate polymer to work with. These difficulties are
enhanced concerning high molecular weight PEO due
to a more complex structural organization of the macromolecule. Indeed, contrary to low molecular weight
PEO obtained from controlled polymerization techniques, high molecular weight PEO are obtained from
condensation of low molecular weight PEO through
multifunctional agent, leading to form both hydrophobic regions and branched structures.
The apparent sensitivity of PEO aqueous solutions
to degradation could be at the origin of some uncertainties that remains concerning the presence of aggregates
and the stability of polymer chains as regard to shear.
Let us focus first on the ability of PEO solution
to form aggregates. Most works refer to the pres-
Rheol Acta (2010) 49:529–540
ence of molecular clusters (or aggregates) in aqueous
PEO solutions, and various interpretations concerning
their origin have been proposed. The aggregation has
been attributed to the presence of impurities in water
(Devanand and Selser 1990), but aggregates have been
shown to spontaneously reformed within 1 day after
filtration of the solution (Polverari and van de Ven
1996; Ho et al. 2003). The role of hydrogen bonds has
also been raised to explain the aggregation phenomena
(Dormidontova 2002), but aggregates seem to be stable
upon heating (Khan 2006; Duval and Sarazin 2003),
a situation for which hydrogen bonds are gradually
broken. Aggregation can be seen as a phase separation
at a microscopic scale. The de Gennes (1991) model,
based on static light scattering data from Polik and
Burchard (1983), proposes that PEO solutions below
70◦ C are phase-separated systems in which aggregates
form a concentrated phase that coexists with a dilute
phase of swollen coils. Such phase separation has been
ascribed to the upper critical solution temperature behavior of PEO solutions. However, a recent smallangle neutron scattering investigation of PEO aqueous
solutions has contradicted this hypothesis (Hammouda
et al. 2004). This study pointed out chain ends effects
on the clustering in PEO solutions with a molecular weight of 4 × 104 g/mol. Despite the fact that
end groups represent only one unit per 1,000 units,
Hammouda et al. (2004) have shown that the ability of
the PEO to aggregate is enhanced in the presence of
nonpolar CH3 groups at both ends of the polymer chain
while it is strongly reduced in the case of chains endcapped by polar OH groups. The contribution of hydrophobic interactions, initially proposed by Polik and
Burchard (1983) and Duval (2000) seems to be determinant when PEO chains are end-capped by nonpolar
groups and would lead to polymer aggregation through
–CH2 –CH2 – groups belonging to the chain and end
groups. However, this interpretation cannot explain
the ability of PEO chain end-capped with OH groups
to form clusters. Several works refer to shear-induced
aggregation in PEO aqueous solutions. According to
Makogon et al. (1988), stirring of high molecular weight
PEO solutions (Mw = 2.4 × 106 g/mol) would favor
steric screening of some ether oxygen atoms, reducing
the affinity of PEO with water and leading to the
dehydration and aggregation of polymer chains without decrease in molecular weight. Rheo-optical measurements have pointed out the trend of micrometric
structures formed under shear to align along the flow
direction (Liberator and McHugh 2005). As stressed by
Hammouda et al. (2004), the aggregation mechanism
depends on experimental conditions and could result from the contribution of various effects, such as
Rheol Acta (2010) 49:529–540
hydrogen bonding, hydrophobic interaction depending
on temperature, and concentration. . .
The stability of PEO chain as regard to flow has also
been the focus of attention in the past decades, but
different interpretations were proposed. It is generally
admitted that the decay of the drag reduction capacity
in high molecular weight (Mw > 106 g/mol) PEO aqueous solutions under elongational flow is due to chain
scission (Minoura et al. 1967; Hunston and Zakin 1978;
Matthys 1991; Sung et al. 2004). A similar effect has
also been reported in high-speed stirring. The number
of bonds broken per chain seems to be independent of
the polymer concentration but increases with increasing
stirring speed (Odell and Keller 1986; D’Almeida and
Dias 1997). The scission rate could be described properly by Jelinek’s or Ovenall’s rate equations (Jellinek
and White 1951; Ovenall et al. 1958) used to characterize ultrasound degradation of polymer solutions. The
scission rate depends also on the nature of the solvent and increases sharply with decreasing the solvent
quality. As observed after ultrasound treatment, chain
scission induced by high-speed stirring is likely to occur
near the center of the macromolecule, leading to the
decrease of the polydispersity index (Buchholz et al.
2004).
According to recent works of Duval et al. (Duval
and Sarazin 2003; Duval and Boué 2007), aggregation
and chain scission are probably combined phenomena.
Indeed, stirring aqueous PEO solutions would lead
first to chain scission and then to the formation of
soluble aggregates. These entities would result from
the intermolecular associations of monomers units previously aligned in the flow direction and stabilized
by dipolar interactions. Aggregates can be dissolved
by the addition of sodium chloride that breaks these
interactions.
The work presented in this paper has been motivated
by the important disparities reported in the literature
in terms of viscosity, intrinsic viscosity, and hydrodynamic radius, respectively, obtained by rheological
or light scattering measurements. For instance, zeroshear viscosities η0 of 4.5 mPa s (Du et al. 2007) and
2 mPa s (Tam and Tiu 1989) are reported for a similar
molecular weight of 4 × 106 g/mol at the concentration
of 8.5 wt.% where polymer coil starts to contact. In
this concentration range, the zero-shear viscosity is
nearly proportional to the concentration, making this
difference in η0 significant. In this work, we present
a thorough investigation of this polymer as a function
of concentration and molecular weight, using rheological techniques. A peculiar attention has been paid to
study the influence of the mechanical history induced
during the dispersion procedure by comparing the rhe-
531
ological properties of polymer solutions prepared by
shaking and by stirring. Dispersion procedures that
govern the mechanical history of samples are generally
summarily treated in the literature. This paper points
out significant changes in the rheological behavior of
polymer solutions according to the mechanical history
of the solutions induced by the dispersion procedures.
Differences are discussed at the molecular scale in
terms of chain and aggregate scission and flow-induced
aggregation.
Experimental section
Materials
Three poly(ethylene oxide) polymers (Sigma-Aldrich)
have been used in this study, each characterized by
an average molecular weight Mw = 5 × 106 g/mol,
106 g/mol, and 3 × 105 g/mol announced by the manufacturer and a polydispersity index of 3.33. Referring
to their expected molecular weight, these polymers are
labeled 5, 1, and 0.3 M, respectively, in the text for more
convenience.
Polymer solutions were obtained by adding the
proper amount of polymer in distilled water and then
by dispersing the polymer at room temperature using
two conventional procedures:
–
–
The softer procedure consists of placing samples
on an SM 30 Edmund Bühler GmbH to-and-fro
motion shaker. A motion of 30 mm amplitude was
applied at the frequency of 175 rpm during a minimum shaking time of 4 days needed to insure a total dispersion of low concentrated solutions, which
can extend to 1 month for the most concentrated
solutions. The dispersion time of 1 month is the upper limit beyond which natural degradation arises,
leading to the decrease of rheological properties
of PEO–water solutions. The homogeneous state
of polymer solutions has been appreciated visually
by checking the uniform natural light scattering of
each sample, and it was confirmed by the reproducibility of rheological measurements for various
samples from the same batch.
The second procedure, broadly used to disperse
polymers, is based on magnetic stirring. Low viscous polymer solutions have been obtained using a
Telesystem 06.40 Thermo Scientific Variomag magnetic stirrer while highly viscous solutions required
an MR Hei-Standard Heidolph Instruments GmbH
and Co. magnetic stirrer. The rotation speed of the
stirring bar, initially immersed in the solvent, was
532
Rheol Acta (2010) 49:529–540
fixed at 500 rpm, and the volume of the stirred sample did not exceed 50 ml to insure an efficient stirring of the whole sample. Polymer solutions have
been continually stirred for 4 days. Samples have
been prepared and stored in darkness to prevent
photooxidation of polymer solutions.
Rheometry
Results and discussion
Influence of the dispersion procedure at various
concentrations for the 5 M PEO
Visual observation and flow curves
The first part of this work is dedicated to the rheological
investigation of the 5 M PEO solutions. In order to
illustrate the significant influence of the dispersion procedure on the rheological properties of the solutions,
two polymer solutions at the concentration of 1.5 wt.%
have been prepared, one by stirring and the other by
shaking. Due to its intermediate level of viscosity, the
1.5 wt.% polymer solution is specially appropriated to
perform both steady shear flow and linear viscoelastic
investigations, and its rheological behavior is representative of the rheological response of PEO solutions. As
a consequence, the 1.5 wt.% polymer solution has been
chosen as a reference for the 5 M PEO. A droplet of
each sample has been placed into the gap of a parallel
plate geometry, and the upper plate has been pulled
up at a constant speed of 5 mm/s, producing an extensional flow. As shown in Fig. 1, the shaken solution
forms a long filament that does not break, even for the
Fig. 1 Basic characterization of elongational properties of
shaken and stirred 5 M PEO solutions at 1.5 wt.%
maximum displacement of the upper geometry, while
the filament for stirred solution is broken for a low
displacement of the geometry. This basic observation
clearly shows that the dispersion procedures used to
disperse high molecular weight PEO in water lead to
the formation of two different final products: a shaken
PEO solution with high elongational properties and a
stirred solution with nearly no elongational property.
To go further in the investigation of the 1.5 wt.% 5M PEO solutions, flow curves of the two systems are
plotted in Fig. 2. Both solutions exhibit a low-shear
Newtonian viscosity, followed by a shear thinning behavior. From a quantitative point of view, the way the
polymer has been dispersed within distilled water has
two major influences on viscous parameters: the zeroshear Newtonian viscosity of the stirred solution of
about 35 Pa s is one decade lower than that of the
shaken solution (η0 ∼ 360 Pa s), while the linear regime
C = 1.5 wt%
10
η (Pa.s)
The linear viscoelastic measurements and steady shear
measurements were carried out using two rotational
rheometers: a controlled strain TA Instruments ARES
rheometer, equipped with a cone and plate geometry
(diameter = 50 mm, cone angle = 1◦ , truncation =
46 μm), and a controlled stress TA Instruments ARG2
rheometer, equipped with a cone and plate geometry
(diameter = 60 mm, cone angle = 0.6◦ , truncation
= 29 μm). The viscosity versus time was measured
at each shear stress or shear rate, and the steadystate viscosity values were determined as the limit, on
long-time scales, of the transient viscosity. To prevent
solvent evaporation during measurements, geometries
were enclosed in a solvent trap which saturates the atmosphere. The lower plate was equipped with a Peltier
thermoelectric device that insures a controlled temperature, fixed at 21 ± 0.1◦ C for this study.
10
10
2
1
0
Shaken solution
Stirred solution
10
-3
10
-2
10
-1
10
0
10
1
10
2
γ (s−1)
Fig. 2 Flow curves of a 1.5 wt.% 5 M PEO solution for which the
polymer has been dispersed by shaking (open symbol) or stirring
(full symbol)
10
10
10
10
10
10
4
3
c
2
4.8
Stirred
1
0
-1
c
-2
-3
c
10
2.7
c
3.3
0.7
-2
10
-1
10
0
c (wt%)
Fig. 3 Zero-shear viscosity of 5 M PEO solutions as a function
of concentration dispersed by shaking (open symbol) and stirring
(full symbol)
In the whole concentration range explored, the zeroshear viscosity of stirred solutions is lower than that of
shaken solutions. Moreover, the two critical concentrations marking the boundaries between concentration
regimes are significantly increased, passing from 0.1
to 0.25 and 0.5 to 0.8 wt.%, respectively, and these
new values are in good agreement with the literature
(Powell and Scharz 1975; Tam and Tiu 1993). From
a phenomenological point of view, the shift of the
concentration regimes towards higher concentrations
associated with a lower value of the viscosity for stirred
solutions can be ascribed to the presence of smaller
molecular entities, i.e., individual polymer chains or aggregates. From a molecular point of view, the decrease
in the size of molecular entities after stirring results
from the mechanical stress of the stirring bar, inducing
either a molecular scission for PEO chains and/or the
breakup of aggregates.
marking the limit of Newtonian behavior is ten times
broader.
This study has been extended to polymer concentrations lying from 0.008 to 5 wt.%. The zero-shear
viscosities of both shaken and stirred polymer solutions have been reported in Fig. 3 as a function of
polymer concentration. For the most concentrated solutions for which the Newtonian behavior is difficult
to access, experimental data have been fitted using the
Bird-Carreau model (Eq. 1) to extrapolate zero-shear
viscosities.
n−1
η = η0 1 + (Kγ̇ )2 2
(1)
Intrinsic viscosity and Huggins coefficient
Three concentration regimes are clearly noticed, each
regime being characterized by power law functions of
η0 with distinct exponents.
ηred = [η] + kH [η]2 c
–
(2)
10000
10000
Shaken
Stirred
3
-1
ηred (cm .g )
For stirring solutions at concentration c∗ < 0.1
wt.%, the power law exponent is found to be equal
to 0.7. In this dilute regime, macromolecules are
supposed to be isolated. The concentration c∗ is
known as the overlap concentration which marks
the transition from a dilute to a semidilute solution
(Doi and Edwards 1986). In the semidilute regime,
0.1 wt.% < c < 0.5 wt.%, η0 increases sharply with
a power law exponent equal to 2.7, corresponding
to the rise of molecular overlaps. The power law
exponent reaches 4.8 in the upper concentrated
regime, for which an entangled network is formed.
For stirred polymer solutions, the power law dependence of η0 is not modified in the diluted
and concentrated regime but the exponent of the
power law of the intermediate regime is 3.3 instead
of 2.7.
8000
6000
4000
3
-1
[η ]shaken = 4000 ± 800 cm .g
kH, shaken= 0.32 ± 0.12
4000
2000
8000
6000
3
-1
[η ]stirred = 900 ± 150 cm .g
kH, stirred = 0.61 ± 0.09
-
–
To attempt to identify the nature of the elementary
objects that contribute to the viscous behavior of PEO
solutions, the reduced viscosity ηred = (η0 − ηw )/ηw c of
both shaken and stirred solutions is plotted in Fig. 4
as a function of concentration c. The reduced viscosity
is defined with η0 , the zero-shear viscosity of polymer
solutions, and ηw = 0.97 mPa s, the Newtonian viscosity
of distilled water at 21◦ C. In the dilute regime, the
reduced viscosity is a linear function of the concentration and could be described by the Huggins (1942)
equation:
-1
η 0 (Pa.s)
10
Shaken
3
10
5
lnη rel /c (cm .g )
10
533
-
Rheol Acta (2010) 49:529–540
2000
0
0.0000
0
0.0005
0.0010
c (g.cm-3)
Fig. 4 Reduced viscosity ηred (open symbols) and inherent viscosity ln(ηrel )/c (full symbols) of shaken and stirred 5 M PEO
solutions as a function of concentration. Intrinsic viscosities are
the extrapolated values to zero concentration of the reduced
viscosity and the inherent viscosity using linear fits
534
where [η] is the intrinsic viscosity and kH the Huggins
coefficient. The intrinsic viscosity, obtained from Fig. 4
by extrapolation of the reduced viscosity at zero concentration, is a unique function of the molecular weight
for a given polymer–solvent pair. Alternatively, [η] can
be obtained by fitting the so-called inherent viscosity,
ηinh = (ln ηrel )/c with the Kraemer equation
ln ηrel
= [η] − kK [η]2 c
c
(3)
where ηrel is the relative viscosity; ηrel = η0 /ηw , and
kK is the Kraemer coefficient. For shaken solutions,
the intrinsic viscosity obtained from both Huggins and
Kraemer representations is about 4,000 ± 800 cm3 g−1
and decreases to 900 ± 150 cm3 g−1 in the case of
stirred solutions. The intrinsic viscosity is related to the
molecular weight M by the Houwink–Mark–Sakurada
(HMS) equation:
[η] = KMα
(4)
where K and α are constants (Flory 1953). For PEO
solutions, HMS constants obtained at T = 25◦ C in a
molecular weight range from 6 × 105 to 106 g/mol are
K = 6.103 × 10−3 cm3 g−1 and α = 0.83 (Khan 2006).
Assuming that HMS constants are not significantly
modified between 25◦ C and 21◦ C and are available for a
broader range of molecular weight, the intrinsic viscosity of about 4,000 ± 800 cm3 g−1 obtained for shaken
solutions corresponds to elementary objects having
an average molecular weight of about (1.02 ± 0.15) ×
107 g/mol. Such big objects, characterized by an average molecular weight significantly higher than that
given by the provider, can be reasonably ascribed to
the presence of aggregates. An elementary observation
confirms this interpretation: during the incorporation of
the polymer in powder form in distilled water, PEO remains at the water surface, forming a hydrated layer of
concentrated and viscous polymer solution. With time,
this layer forms a macroscopic aggregate that vanishes
by diffusion mechanisms. This observation stresses the
spontaneous tendency of PEO chains to aggregate in
distilled water at room temperature (T ∼ 18◦ C). The
dispersion of polymer under shaking is mainly due to
the diffusion of water molecules within the hydrated
polymer layer, and the mechanical energy brought during shaking is significantly lower compared to the stirring process. As a consequence, the presence of PEO
clusters is highly expected under shaking.
On the contrary, for stirred solutions, the intrinsic
viscosity of about 900 ± 150 cm3 g−1 corresponds to
elementary objects having an average molecular weight
of about (1.69 ± 0.2) × 106 g/mol, which is well below
the theoretical value and could be due to chain scis-
Rheol Acta (2010) 49:529–540
sion after stirring. This interpretation is consistent with
the spectacular decrease of the elongational behavior
pointed out for stirred solutions in Fig. 1. Indeed, the
elongation at break is known to significantly decrease
with decreasing molecular weight. In the case of polymer solutions, elongational properties are governed
by dispersed chains, and the contribution of possible
aggregates is negligible. Consequently, differences in
the elongational properties between shaken and stirred
polymer solution can be clearly understood with a decrease of PEO molecular weight, confirming our interpretation of polymer scission during stirring.
In order to have physicochemical insight into the
microstructure of this complex polymeric system, we
have considered the Huggins coefficient kH obtained
from the slopes of ηred in Fig. 4. The value of kH
that depends on the solvent quality is linked to the
second virial coefficient A2 (Yamakawa 1961) and is
generally considered as a parameter that quantifies the
interactions between two molecules also called pair
interactions. For shaken solutions, kH,shaken = 0.32 ±
0.12 while kH,stirred = 0.61 ± 0.09 in the case of stirred
solutions. The lower value of kH,shaken suggests that
aggregates interact weakly and can be considered as
isolated entities in dilute solutions. On the contrary,
the higher value of kH,stirred implies that PEO chains
interact strongly with neighboring chains. The more
associative character of stirred polymer solutions could
be assigned to chain scissions. Indeed, polymer scission could break C–C or C–O bonds with nearly the
same probability since their bond energy of 83 and
81 kcal mol−1 , respectively, are very close (Cottrell
1958). The specific interactions of the oxygen lone pair
with the hydrogen of water, which bears a partial positive charge, may however influence the predominant
bond breaking. If the breaking occurs at the C–C bonds,
it will result in the formation of two R − O − CH•2
where R designates the polymer chains. If the breaking
occurs at the C–O bonds, it will result in the formation
of two radicals R − O − CH2 − CH•2 and R–O• . The
stability of carbon radicals decreases as follows: R −
O• > R − O − CH•2 > R − O − CH2 − CH•2 . As a C–O
breaking provides both the more stable and less stable
radicals while the C–C bond breaking provides two
radicals of intermediary stability, both types of breakings are realistic. Chains end-capped by free radicals
can recombine, giving a variety of polymer chains but
involve no modification of the end-group nature.
Broken bonds giving rise to −CH•2 and –O• radicals
could lead to the formation of –CH3 and –OH end
groups, respectively, through H• transfer mechanism.
This H• transfer may proceed from a transfer to polymer; the resulting radical may recombine with the other
macromolecular chain fragments, leading to a branched
PEO. When the H• transfer is a transfer solvent, it
generates a very aggressive hydrophilic OH• radical
that can attack the polymer chain.
The formation of –OH end groups through polymer scission would not modify interactions with the
solvent. On the contrary, the presence of hydrophobic –CH3 , –CH=CH2 , or –CH2 –CH3 , depending on
where the polymer scission takes place in the ethylene
groups, favors associative interactions between broken
polymer chains. Such hydrophobic interactions may
contribute to the high pair interactions measured for
stirred PEO solutions. Despite their small fraction, end
groups are found to play a dominant role in PEO
solutions (Dormidontova 2004; Hammouda et al. 2004).
The density of end groups, i.e., their contribution in
the rheological response of polymer solutions, increases
when decreasing the molecular weight. Hammouda
et al. 2004 have suggested that hydrophobic interactions
between end groups and ethylene groups of the chains
are at the origin of clustering in PEO solutions. We
cannot exclude the formation of small aggregates in
dilute solutions dispersed by stirring, but their rheological signature is not noticeable. So, the schematic
insight into PEO chain organization in dilute solutions
is that shaking dispersion procedure could favor the
formation of aggregates of nondegraded chains while
stirring dispersion prevents their formation and induces
chain scission.
Linear viscoelastic measurements and relaxation
time dynamics
In this dilute regime, where polymer chains are not entangled, the elongational stress that induces chain scission during stirring is most likely due to hydrodynamic
forces transmitted by the suspending medium through
friction. What happens in the concentrated regimes for
which polymer chains are entangled? To answer this
question, the rheological investigation has been completed by linear viscoelastic measurements for polymer
solutions at a concentration higher than 0.5 wt.%, for
which a viscoelastic response can be measured. Figure 5
shows the frequency dependence of the storage modulus G and loss modulus G for the reference solutions
at 1.5 wt.%. For stirred solutions, G and G moduli
are, respectively, proportional to ω2 and ω1 at low
frequencies, corresponding to the terminal portion of
the curves, while both moduli increase slowly at higher
explored frequencies. Such a linear viscoelastic behavior is observed in dense macromolecular systems having
a broad relaxation time distribution, generally ascribed
to the molecular polydispersity, and for which long-
10
10
535
2
1
Shaken
G' ; G'' (Pa)
Rheol Acta (2010) 49:529–540
10
0
0.6
10
10
10
-1
ω
1
Stirred
1
ω
2
ω
-2
G'
G''
-3
10
-4
10
-3
10
-2
10
-1
10
0
10
1
10
2
ω (rad/s)
Fig. 5 Storage modulus G and loss modulus G versus frequency of 1.5 wt.% 5 M PEO solutions dispersed by shaking
(open symbols) and stirring (full symbols). Continuous lines
through data represent the fit using the generalized Maxwell
model. Dot lines are added to account for the power law dependence of shaken solutions at low frequencies
time dynamics, characterized by the terminal zone, is
governed by reptation (de Gennes 1979).
For shaken PEO solution, it is first worth noting
that the viscoelastic moduli are higher than those reported for stirred solutions. This observation is valid
in the whole explored frequency range. A significant
departure of G and G moduli from the terminal zone
is observed at lower frequencies. Indeed, the pulsation dependence of viscoelastic moduli follows a power
law of ω1 and ω0.6 , respectively, in this zone. Such a
weaker power law dependence of viscoelastic moduli
in the terminal zone has already been noticed for soft
microgel suspensions of commercial associative guar
gum (Aubry et al. 2002) and confirms the presence of
aggregates in shaken solutions as suggested previously
in this work.
The linear viscoelastic behavior can be expressed as
a function of a discrete spectrum using the generalized
Maxwell model:
G (ω) =
m
Gi
i=1
G (ω) =
m
i=1
(ωλi )2
1 + (ωλi )2
(5)
ωλi
1 + (ωλi )2
(6)
Gi
where Gi is the elastic modulus corresponding to a
relaxation time λi . The fit of G and G data using
Eqs. 5 and 6 has been performed in order to monitor
the concentration dependence of the longest relaxation time, ascribed to the slow dynamics of the longest
G'
G''
2
λ s (s)
c
10
Shaken
3.8
1
c
2.5
Stirred
10
0
10
0
10
1
c (wt%)
Fig. 6 Concentration dependence of the longest relaxation time,
λs , for shaken (open symbols) and stirred (full symbols) 5 M PEO
solutions obtained by fitting linear viscoelastic moduli with the
generalized Maxwell model
molecular entities that are individual chains or aggregates. These slower relaxation time dynamics, noticed
λs , are reported as a function of concentration in Fig. 6.
Both G and G moduli have been fitted separately and
give very close values of λs for each concentration. Error bars reported in Fig. 6 correspond to the uncertainty
on relaxation times obtained from the fitting.
Let us consider first the shaken solutions. According
to the previous discussion, the high values of λs for
shaken solutions are probably associated with the relaxation dynamics of aggregates while relaxation dynamics
of long chains dispersed in the suspending medium
are represented by the distribution of faster relaxation times. In the concentration regime ranging from
0.5 wt.% to a critical concentration of about 1.25 wt.%,
marking roughly the beginning of the entangled regime,
the relaxation time increases with the concentration
according to a power law with an exponent of about
3.8. It is worthwhile to note that the concentration
dependence of λs is much more important than that expected, with a power law of 0.5 generally observed for
well-dispersed chains in the semidilute nonentangled
regime. The discrepancy between exponents could be
explained by the presence of aggregates for shaken solutions: for polymer solutions of individualized chains,
the increase of the relaxation dynamics with increasing
polymer concentration is due to an increase of constraints induced by neighboring chains. For aggregates,
their relaxation dynamics increases with increasing the
aggregate size. By adding polymer, the relaxation dynamics of aggregates is more efficiently hindered than
that of single molecules.
In addition, a direct relation between zero-shear
viscosity and relaxation time can be obtained by com-
Rheol Acta (2010) 49:529–540
paring their concentration dependence. Below the critical concentration regime, the zero-shear viscosity for
shaken solutions is found to be proportional to the
ratio λi /c. In the entangled regime c > 1.25 wt.%, λs
still increases with concentration but seems to level
off in the highest concentration regime. It is worth
mentioning that λs values in the highest concentration
regime are more difficult to obtain, and this result will
not be commented.
Shaken solutions can be reasonably considered as
suspensions of polymeric aggregates dispersed in a water solution of long polymer chains. Such coexistence
of a concentrated phase dispersed in a more dilute
phase has been proposed by de Gennes (1991) for PEO
solutions in water and can be evidenced by the turbid
character of polymer solution for concentrations higher
than ∼1 wt.%.
What about the slow relaxation time dynamics for
stirred solutions? With increasing concentration in the
lowest concentration regime 0.5 wt.% < c < 1.25 wt.%,
λs increases with an exponent of the power law of about
2.5, and the zero-shear viscosity is found to be proportional to the product λs · c. Then, λs reaches a maximum
around the critical concentration and finally tends to a
nearly constant value at high polymer concentrations.
The exponent of the power law of λs is lower than that
obtained for shaken solution, but it is still higher than
the expected value. This point will be discussed further
in the article.
Discussion of results at the molecular scale is more
delicate for concentrated solutions prepared by stirring.
For this purpose, two stirred solutions at 1.5 wt.%
6000
Dilution from c = 1.5wt%
Dilution from c = 3.6wt%
-1
10
3
4000
3
[η]1.5= 3620 cm /g
3
10
η red (cm .g )
536
2000
3
[η]3.6= 1800 cm /g
0
0.0000
0.0005
c
0.0010
(g.cm-3)
Fig. 7 Reduced viscosity of stirred 5 M PEO solutions obtained
from the dilution of a 1.5 wt.% mother solution and a 3.6 wt.%
mother solution as a function of concentration. Intrinsic viscosities are the extrapolated values to zero concentration of the
reduced viscosity using linear fits
–
537
As pointed out by Pang and Englezos (2002) for
phase-separated PEO, aggregates are very sensitive
to shear and can breakup under flow.
The molecular dynamics seems to be governed by
the competition between the hindering effects induced
by the growth rate of aggregate size and the speed up
effect induced by chain and aggregate scissions.
An increase of the polymer concentration below the
critical concentration of 1.25 wt.% seems to favor the
hindering effect with a gradual growth of aggregates.
The lower exponent of the power law of λs between
shaken and stirred solutions may be ascribed to the
weaker increase of the growth rate of aggregate size
induced by shorter polymer molecules compared to
nondegraded chains for shaken solutions.
As polymer concentration increases above the critical concentration of 1.25 wt.%:
–
Entanglements induces additional stresses that favors the formation of shorter chains with a more
randomized scission, probably between knots of the
entangled network that leads to increase of the
polydispersity index (Odell et al. 1992).
Frictions between aggregates and hydrodynamic
interactions with the more viscous suspending
medium increase, favoring the breakup of aggregates.
–
In this latter case, the breakup of aggregates is probably more important than the growth rate of aggregates, leading to the formation of smaller aggregates at
high concentrations.
10
10
10
10
3
Shaken
Stirred
2
5
c
1
0
0
corresponding to the highest value of λs and 3.6 wt.%
for which λs is low have been diluted in the concentration range c < 0.01 wt.% in order to estimate their
intrinsic viscosity and the size of individual objects
formed in the solution. Figure 7 shows the reduced
viscosity of stirred solutions obtained by the dilution
of two mother stirred solutions at 1.5 and 3.6 wt.%
as a function of concentration. The intrinsic viscosity
[η]1.5 ∼ 3,620 cm3 g−1 of solutions obtained by dilutions
of the 1.5 wt.% is significantly higher than [η]3.6 ∼
1,800 cm3 g−1 obtained by dilutions of the 3.6 wt.%.
According to Eq. 4, stirred solutions at 1.5 wt.% would
contain polymeric objects with an average molecular
weight of about 9.0 × 106 g/mol, while at 3.6 wt.%, the
average molecular weight is about 3.9 × 106 g/mol.
Stirring has been shown to reduce the average molecular weight of PEO. Moreover, in the semidilute and
concentrated regime, PEO chains are likely to form aggregates with increasing concentration as chain contact
increases considerably. This aggregation mechanism is
macroscopically evidenced by a more turbid sample,
also observed for shaken solutions. Consequently, both
molecular weights have to be considered as the signature of an aggregation state of short polymer chains
for which the molecular weight determination is difficult. It has to be stressed that [η]1.5 > [η]3.6 , and both
are significantly higher than that obtained from dilute
solutions in Fig. 4, [η] ∼ 900 cm3 g−1 . This observation
shows that aggregates formed in the semidilute regime
are bigger that those obtained in the concentrated
regime, themselves bigger than probably individualized
and degraded PEO chains. Moreover, this result suggests that big aggregates formed under stirring as concentration increases are stable in water after dilution
and appear consequently as insoluble entities, strengthening the interpretation of the hydrophobic character
of PEO aggregates. Indeed, the hydrophile/lipophile
balance that quantifies the hydrophilic character of
PEO chains decreases with both increasing the concentration (Kim and Cao 1993) and decreasing the molecular weight (Cao and Kim 1994) due to hydrophobic end
effects.
At this point of the study, we can propose a
schematic insight into PEO chain organization in semidilute and concentrated solutions dispersed by stirring. In these concentration regimes, stirred solutions
contain aggregates for which the effect of stirring is
twofold:
η (Pa.s)
Rheol Acta (2010) 49:529–540
10
10
10
-1
2.4
c
-2
0.6
c
-3
10
-1
10
0
10
1
c (wt%)
–
It favors the formation of shorter chains having an
enhanced hydrophobic character that contributes
to aggregation.
Fig. 8 Zero-shear viscosity of 1-M (square symbols) and 0.3 M
(triangle symbols) PEO solutions as a function of concentration
and dispersed by shaking (open symbol) and stirring (full symbol)
538
Influence of the molecular weight
ηred (cm3.g-1)
1500
In order to complete the physical insight into the microstructure of PEO solutions, rheological measurements have been extended to two lower molecular
weight polymers with an expected value of 106 and
3 × 105 g/mol, labeled 1 and 0.3 M. Figure 8 shows the
zero-shear viscosity curves for both molecular weight
polymer solutions prepared by shaking and stirring as a
function of polymer concentration.
For 1 M polymer solutions in the dilute regime c <
0.5 wt.%, η0 does not depend on the dispersion procedure. This result suggests that hydrodynamic forces
transmitted by the suspending medium are not strong
enough to produce chain scission of PEO solutions with
molecular weight lower than 1 M. Indeed, the critical
elongational strain rate ε̇ f required for chain scission of
high molecular weight PEO in dilute regime is found
to be proportional to Mw−2.25 (Islam et al. 2004). By
decreasing Mw below 1 M, ε̇ f cannot be achieved under
stirring at 500 rpm, and PEO chains are stretched but
do not break.
In the semidilute regime, the zero-shear viscosity of
stirred solutions is slightly lower than that of shaken
solutions, and this difference is more pronounced in
the concentrated regime. However, the gap between
zero-shear viscosities of shaken and stirred solutions
is lower than that observed previously for the 5 M
molecular weight polymer. As the concentration increases, polymer chains undergo additional stress due
to entanglements that could be responsible for either
chain and/or aggregate scission. This additional stress
enhances as molecular weight increases.
For the 0.3 M PEO solutions, no difference in η0
is noticeable between stirred and shaken solutions in
the whole range of concentration explored. Indeed, it
has been reported that the ability of PEO chains to
form aggregates vanishes below the critical molecular
weight of about 6 × 105 g/mol, and macromolecules
behave as flexible chains in a good solvent (Rangelov
and Brown 2000). For such low molecular weights in
the concentrated regime, elongational stresses that are
strengthened by local stresses of neighbor chains are
not strong enough to induce chain scission. It has to be
noticed that concentration dependence of η0 seems to
be independent on both the molecular weight and the
dispersion procedure and characterized by a power law
with an exponent of about 0.6 in the dilute regime and
5 in the concentrated regime.
The reduced viscosity ηred for both molecular weight
polymer solutions obtained by shaking and stirring is
plotted in Fig. 9 as a function of concentration. Con-
Rheol Acta (2010) 49:529–540
Shaken
Stirred
1000
[η]1M=820 g.cm
-3
500
[η]0.3M=250 g.cm
0
0.000
0.002
-3
0.004
0.006
0.008
-3
c (g.cm )
Fig. 9 Reduced viscosity of 1 M (square symbols) and 0.3-M
(triangle symbols) PEO solutions as a function of concentration
and dispersed by shaking (open symbol) and stirring (full symbol)
trary to results obtained for 5 M PEO, ηred for stirred
and shaken solutions is very close for the 1 M polymer, especially at low concentrations, and superimpose
for the 0.3 M PEO, forming a master curve. The intrinsic viscosities [η]1 M ∼ 820 cm3 g−1 and [η]0.3 M ∼
250 cm3 g−1 extrapolated from Fig. 9 correspond, respectively, to objects with an average molecular weight
of about 1.5 × 106 and 3.6 × 105 g/mol. Both values are
close to the expected molecular weight, confirming that
polymer scission induced by stirring dilute water PEO
solutions at 500 rpm occurs only for molecular weight
higher than about 106 g/mol.
Conclusion
Rheological properties of PEO aqueous solutions have
been characterized as a function of the polymer concentration and molecular weight for two dispersion
processes: stirring and shaking. The whole set of data
reported in this paper clearly shows that shaking favors
aggregation and stirring favors chain scission, depending on concentration and molecular weight.
For Mw ≤ 3 × 105 g/mol, neither chain scission nor
aggregation is noticed using rheological investigation
techniques. As a consequence, the dispersion process
does not modify the rheological properties of low molecular weight PEO solutions. However, chain scission
and aggregation cannot be excluded.
For Mw ≥ 3 × 105 g/mol, differences in composition
appear in PEO solutions according to the concentration
and the dispersion procedure. In the dilute regime,
moderated stirring is able to break PEO long chains
via hydrodynamic forces while shaking leads to the
Rheol Acta (2010) 49:529–540
formation of aggregates. The ability of polymer scission
induced by the dispersion procedure increases with
increasing Mw . In the concentrated regime, aggregation
of PEO chains is favored by increasing the concentration. Hydrodynamic forces combined to additional
stresses due to entanglement enhance the breakup of
polymer chains and aggregates.
These results are of potential practical importance as
regards to industrial applications involving frequently
turbulent flow processes. In particular, it stresses that
the mechanical history of PEO solutions leads to the
formation of polymer solutions having differences in
their compositions (aggregates, chain scission) and in
their rheological properties. The general rule that recommends the use of high molecular weight to provide
high viscous and viscoelastic solutions is still available
but should be balanced by considering the dispersion
process.
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Electrochimica Acta 55 (2010) 5186–5194
Contents lists available at ScienceDirect
Electrochimica Acta
journal homepage: www.elsevier.com/locate/electacta
Nanocomposite polymer electrolyte based on whisker or microfibrils
polyoxyethylene nanocomposites
Fannie Alloin b,∗ , Alessandra D’Aprea a,b,c , Nadia El Kissi a , Alain Dufresne c , Frédéric Bossard a
a
Laboratoire de Rhéologie, Grenoble-INP-UJF, UMR 5520, BP 53, 38041 Grenoble Cedex 9, France
LEPMI, Laboratoire d’Electrochimie et de Physicochimie des Matériaux et des Interfaces, Grenoble-INP-UJF-CNRS, UMR 5631, BP 75, 38041 Grenoble Cedex 9, France
c
Ecole Internationale du Papier, de la communication imprimée et des Biomatériaux, PAGORA- Grenoble-INP, BP 65, 38402 Saint Martin d’Hères Cedex, France
b
a r t i c l e
i n f o
Article history:
Received 22 February 2010
Received in revised form 8 April 2010
Accepted 8 April 2010
Available online 2 May 2010
Keywords:
Electrolyte
PEO
Nanocomposite
Lithium battery
Mechanical property
a b s t r a c t
Nanocomposite polymer electrolytes composed of high molecular weight poly(oxyethylene) PEO as a
matrix, LiTFSI as lithium salt and ramie, cotton and sisal whiskers with high aspect ratio and sisal
microfibrils (MF), as reinforcing phase were prepared by casting-evaporation. The morphology of the
composite electrolytes was investigated by scanning electron microscopy and their thermal behavior
(characteristic temperatures, degradation temperature) were investigated by thermogravimetric analysis
and differential scanning calorimetry.
Nanocomposite electrolytes based on PEO reinforced by whiskers and MF sisal exhibited very high
mechanical performance with a storage modulus of 160 MPa at high temperature. A weak decrease of the
ionic conductivity was observed with the incorporation of 6 wt% of whiskers. The addition of microfibrils
involved a larger decrease of the conductivity. This difference may be associated to the more restricted
PEO mobility due to the addition of entangled nanofibers.
© 2010 Elsevier Ltd. All rights reserved.
1. Introduction
Polymer-based ion conducting materials have generated
remarkable interest in the field of lithium batteries thanks to their
application as electrolytes since Armand [1] proposed the use of
poly(oxyethylene), PEO and lithium salt as solid polymer electrolyte. Conductivity of electrolytes based on PEO, a semicrystalline
polymer, strongly depends on the crystalline phase proportion,
which is usually considered as poorly conductive [2], whereas in
the amorphous phase, the ionic mobility is assisted by the polymer
motion. The ionic conduction, below the melting point, depends
strongly on the thermal history of the sample. The melting point
of the polymer may be decreased by using a “plasticizer” salt, such
as LiTFSI [3], lithium bis(trifluoromethane sulfonyl)imide. Furthermore, due to the strong electron-withdrawing of SO2 CF3 groups at
both sides of the imide anion, LiTFSI presents a high salt dissociation
level even in a low dielectric medium such as polyether.
According to the salt used and its concentration, the crystallization kinetic of the PEO electrolyte can be very low, with a complete
crystallization, at room temperature, after several days. Crystallization causes a considerable drop of ionic conductivity in PEO based
electrolytes [4,5].
∗ Corresponding author. Tel.: +33 4 76 82 65 60; fax: +33 4 76 82 66 70.
E-mail address: [email protected] (F. Alloin).
0013-4686/$ – see front matter © 2010 Elsevier Ltd. All rights reserved.
doi:10.1016/j.electacta.2010.04.034
Moreover, for safety and performance reasons, polymer electrolyte must exhibit, in addition to high conductivity and a wide
electrochemical stability window, high thermal and mechanical
performances. Since the original work of Weston and Steel [6],
who reported the improvement of polymer electrolyte conductivity
and mechanical stability by adding ␣-Al2 O3 particles, nanocomposite polymer electrolytes have been extensively studied [7,8].
The understanding of the impact of the inorganic filler on conduction, thermal, mechanical and electrochemical properties of
polymer electrolytes are still in progress. The best performances
were obtained with Al2 O3 [8,9] and TiO2 [7]. The increase in conductivity was found to be significant below the melting point of
PEO. This improvement was ascribed to the highest amorphous
state of the electrolyte, due to the decrease of the crystallization
kinetic induced by the fillers. In PEO/LiTFSI complexes, an amorphous structure, stable for several months, can be obtained with
the incorporation of small amounts of fillers [10]. Using Li NMR
investigation, Dai et al. [11] found that the addition of nanometric
Al2 O3 to PEO/LiI electrolyte suppressed the formation of crystalline
phases.
Cellulosic rigid rods whiskers extracted from tunicate, a sea
animal [12,13] were used as mechanical reinforcing phase in saltfree PEO based composites and composite electrolytes. Spectacular
improvement of the tensile modulus especially above the melting
temperature, with a value of 20 MPa, was observed with only 3 wt%
cellulosic fillers. This elevated reinforcing effect was ascribed to (i)
the high aspect ratio of the fillers and (ii) the formation of a per-
F. Alloin et al. / Electrochimica Acta 55 (2010) 5186–5194
colated cellulosic network within the polymeric matrix above the
percolation threshold. This behavior is due to strong interactions
between whiskers by hydrogen bonds that form a rigid network. A
small decrease of the ionic conduction was observed and associated
to the decrease of the PEO mobility due to PEO/whiskers interaction. However, this extraordinary reinforcement was obtained
with a model organic filler, i.e. whiskers extracted from tunicate.
These whiskers cannot be employed easily because of their hard
supplying and high cost. In a previous study [14], a remarkable
improvement of the tensile modulus with organic whiskers issued
from ramie plant, easier to find, on the PEO matrix was obtained.
The present work reports the use of sisal, ramie and cotton
whiskers and sisal microfibrils as organic fillers in PEO matrix
and PEO–LiTFSI polymer electrolytes. Different sources of cellulose plants were employed in order to evaluate the influence of
the charge density and the aspect ratio of the fillers on composite
polymers and composite electrolyte properties. Furthermore, the
effect of both sisal whiskers and microfibrils has been studied.
2. Experimental
2.1. Composite films
2.1.1. Polymeric matrix
Poly(oxyethylene), PEO, with high molecular weight
(Mw = 5 × 106 g mol−1 ) was purchased as a white powder
from Aldrich. The polymer was used as received. The lithium
bis(trifluoromethane sulfonyl imide) LiTFSI from Fluka was dried
under vacuum for 48 h at 130 ◦ C and then stored in glove box.
2.1.2. Cellulose nanocrystals
2.1.2.1. Sisal whiskers. Cellulose whiskers were extracted from
sisal plant originating from Northeast Brazil and purchased from
Mariana (Minas Gerais, Brazil). Native sisal leaves were cut with
a 300 W mixer until a fine fibrous powder was obtained that
was subsequently washed four times in boiling 2 wt% aqueous
NaOH for 4 h under mechanical stirring. The material was filtered and rinsed with distilled water between each treatment
step. A bleaching treatment at 80 ◦ C was used to bleach the sisal
fillers. The bleaching solution contained equal parts of aqueous
chlorite and acetate buffer. The sisal content was approximately
5 wt% and the bleaching step was repeated twice. The fillers
were filtered and rinsed with distilled water between each treatment step and then dried for 24 h at 40 ◦ C in a convection
oven. The dried fillers were ground a second time to a fine
powder using a 300 W mixer and dispersed in 65 wt% sulfuric acid in water (4 wt% sisal). This suspension was diluted and
washed by successive centrifugations at 10 rpm and 10 ◦ C. Dialysis with distilled water and sonication were done before storing
the whiskers in the refrigerator with several drops of chloroform.
2.1.2.2. Cotton whiskers. Suspension at 8 wt% of cotton was washed
with a solution of 64 wt% of H2 SO4 during 45 min at 45 ◦ C. This
suspension was diluted with distilled water and centrifuged at
11,000 rpm several times. Dialysis with distilled water, washing
with resin and sonication were done before storing the whiskers
suspension in the refrigerator with several drops of chloroform.
The length L and cross-section d of these nanoparticles were
estimated, by TEM analysis, at about 165 ± 34 nm and 13 ± 1 nm,
respectively, after 200 measurements. The average aspect ratio L/d
and the specific area of these whiskers were calculated to be close
to 13 ± 3 m2 g−1 and 205 ± 25 m2 g−1 , respectively.
2.1.2.3. Ramie whiskers. The treatments of Ramie plants in order
to obtain whiskers was described elsewhere [15,16]. Briefly, ramie
5187
fillers were cut into small pieces and treated with 2 wt% NaOH at
80 ◦ C for 2 h to remove residual additives. The purified ramie fillers
were submitted to acid hydrolysis with a 65 wt% H2 SO4 solution
at 55 ◦ C for 30 min and under continuous stirring. The suspension
was washed with water by centrifugation and dialyzed to neutrality
against deionized water. The obtained suspension was homogenized using an Ultra Turrax T25 homogenizer and then filtered to
remove unhydrolyzed fillers. The whiskers were stored in refrigerator with several drops of chloroform.
2.1.2.4. Sisal microfibrils (MF). Sisal microfibrils were prepared
from sisal fillers as described elsewhere [13]. Briefly, a 2.0 wt% solution of bleached sisal fillers was pumped through a microfluidizer
processor, Model M-110 EH-30. The slurry was passed through
the valves that applied a high pressure. Size reduction of products
occurs into Interaction Chamber (IXC) using cellules of different
sizes (400 ␮m and 200 ␮m). Pumping cycles were varied in order
to optimize the fibrillation process. The cross-section and specific
area of MF were determined and were close to 36 ± 12 nm and
74 ± 7 m2 g−1 , respectively.
2.1.3. Films processing
2.1.3.1. Whiskers as reinforcement. The whiskers suspension was
sonicated, in order to obtain a stable suspension, 5 min in ice
bath. The desired amount of sisal whiskers (aqueous suspension)
was added to the PEO previously dispersed in a few drops of
methanol. The resulting suspension was protected against light by
an aluminum foil and was weakly stirred for 4 days at room temperature in order to avoid PEO degradation [17]. The suspension
was degassed under vacuum in order to remove the remaining air
and cast into Teflon plates under argon at 40 ◦ C for 3 weeks. Films
were dried under vacuum to eliminate the remaining water for 1
week with a step of 5 ◦ C every day from 40 ◦ C to 75 ◦ C. The samples
were then stored in glove box.
Desired amount of salt LiTFSI was dissolved in a few milliliters of
acetonitrile and introduced by swelling the nanocomposite films in
glove box. This procedure was selected since the whiskers dispersion stability is ensured by electrostatic repulsions. The addition
of salt induces the charge screening, thus the electrostatic repulsion between whiskers decrease and the whiskers self-aggregate
and settle down, leading to a phase separation of the suspension.
Finally, nanocomposite electrolytes were dried under vacuum at
60 ◦ C for 72 h and stored in glove box. The final films were around
200–300 ␮m thick.
2.1.3.2. Sisal MF as reinforcement. As Sisal MF aqueous suspensions
are inherently stable no sonication step was necessary. The desired
amount of MF (aqueous suspension) was added to PEO dispersed
in a few ml of methanol. For the electrolyte elaboration, the same
procedure was used and the salt was added during the preparation
of the composite film.
Other steps used for the preparation of the composite salt-free
samples and electrolytes were similar to those previously described
for whisker-based samples.
2.2. Measurements
2.2.1. Microscopies
Scanning electron microscopy (SEM) was used to investigate the
morphology of the nanocomposite films using a LEO S440 SEM
instruments. The specimens were frozen under liquid nitrogen,
fractured, mounted, coated with graphite and observed using an
accelerating voltage of 10 kV.
Transmission electron microscopy (TEM) observations were
made with a Philips CM200 electron microscope. A droplet of a
dilute suspension of cellulose nanoparticles was deposited and
5188
Table 1
Dimensions and charge surface density of sisal, ramie, cotton whiskers and sisal MF.
Sisal whiskers [14]
Ramie whiskers [15]
Cotton whiskersb
Sisal MFb
a
b
Length L (nm)
Cross-section d (nm)
215 ± 67
200 ± 78
165 ± 34
–
5
6–8
13
36
±
±
±
±
1.5
1
1
12
L/d (length/cross-section)
43 ± 26
28 ± 12
13 ± 3
–
Specific area (m2 g−1 )
600
380
205
74
±
±
±
±
48
38
25
7
Charge surface densitya (e/nm2 )
0.026 ± 0.001
0.021 ± 0.001
0.029 ± 0.001
–
e the electron charge, 1.6 × 10−19 C.
Determined in this work.
allowed to dry on a carbon-coated grid. The accelerating voltage
was 80 kV.
2.2.2. Differential scanning calorimetry
Differential scanning calorimetry (DSC) was performed using a
TA Instrument DSC DSC2920 CE. Standard mode was used. Around
10 mg of samples were placed in a DSC cell in glove box. Each sample was heated from −100 ◦ C to 100 ◦ C at a temperature ramp of
10 ◦ C/min and kept at 100 ◦ C for 5 min to insure complete melting
of composite films and electrolytes. Then it was cooled down to
0 ◦ C at a cooling rate of 10 ◦ C/min. The temperature and enthalpy
of melting were determined during the first DSC scan in order to
analyze samples at equilibrium.
The melting temperature, Tm , and the crystallization temperature, Tc , were taken at the onset of the peaks corresponding to the
melting endotherm and the crystallization exotherm, respectively.
2.2.3. Thermogravimetric analysis
Thermogravimetric measurements were carried out with a Netzsch STA409 thermal analyzer. Around 10 mg of samples were
heated from room temperature up to 550 ◦ C at 10 ◦ C/min under
nitrogen flow. A thermobalance determined the sample weight loss
under non-isothermal temperature. The degradation temperature
was taken at the onset of the weight loss.
2.2.4. Dynamic mechanical analysis
Dynamic mechanical analysis (DMA) measurements were carried out with a spectrometer DMA Q800 from TA Instrument
working in the tensile mode. The strain magnitude was fixed at
0.05%. This value ensures that tests were made in the linear viscoelastic domain.
The samples were thin rectangular strips with dimension of
about 20 mm × 7 mm × 0.2 mm. Measurements show the storage
tensile modulus vs temperature.
2.2.5. Conductivity measurement
Ionic conductivity was measured by impedance spectroscopy
using a HP4192A impedance analyzer, over the frequency range
5–13 MHz. The samples were placed between two stainless-steel
blocking electrodes under argon. The temperature sweep test was
conducted from 25 ◦ C to 75 ◦ C. The temperature was equilibrated
1 h every 5 ◦ C between each measurement.
Abbreviation. The lithium salt content in the film was classically referring to the number n = O/Li, which corresponds to the
molar ratio oxyethylene unit/lithium. PEO based electrolyte will be
labeled as PEO-O/Li = X + Y wt% filler name, X being the O/Li ratio
and Y the whiskers content.
3. Results and discussion
3.1. Morphology
3.1.1. Salt-free PEO nanocomposites
Salt-free PEO nanocomposites were obtained using different cellulose whiskers (cotton, ramie and sisal plants) as well as sisal
MF. The whiskers and MF had different dimensions, thus different
aspect ratios, specific areas and charge surfaces (Table 1). Indeed,
the stiffness and the aspect ratio of the whiskers have been shown
to depend on the type of plant and on the degree of crystallinity of
the fibers and thus nanocomposites with different properties may
be obtained.
The length of MF was difficult to determine and can be considered as very important compared to its cross-section. The sisal
whiskers present the highest specific area, whereas the sisal MF
exhibits the lowest one. The aspect ratio, decisive for mechanical
property, is between 43 and 13 for the different whiskers, thus a
large difference in mechanical property is expected.
The morphology of the cellulosic fillers/PEO composites was
characterized by SEM. Fig. 1 shows the morphology of both cotton whiskers/PEO composites and MF sisal/PEO composites. For all
composite samples, a homogeneous dispersion of white dots was
observed and has been associated with the presence of whiskers
or microfibrils. These white dots do not correspond directly to isolated particles. Indeed, the particle dimensions were too small to be
observed at this scale. The white dots result from electrical charge
Fig. 1. Scanning electron micrographs of the cryo-fractured surface of (a) PEO + 6 wt% cotton whiskers nanocomposite film and (b) PEO + 6 wt% sisal MF nanocomposite film.
5189
Table 2
Thermal characteristics of PEO based nanocomposite films reinforced with sisal
microfibrils (MF) and sisal (WS), ramie (WR) and cotton (WC) whiskers obtained
from DSC curves: crystallization temperature (Tc ), melting temperature (Tm ), associated enthalpy of melting per gram of PEO (Hm-PEO ) and degree of crystallinity
measured during the melting (m ). The degradation temperature Tonset corresponds
to the beginning of the degradation process performed in helium.
PEO
PEO + 6% WS
PEO + 6% WC
PEO + 6% WR
PEO + 6% MF
Tm (◦ C)
Hm-PEO (J/g)
m
Tc (◦ C)
Tonset He (◦ C)
56
56
57
57
53
170
165
162
170
167
0.81
0.77
0.74
0.8
0.79
51
48
50
50
44
350
260
260
357
350
The degree of crystallinity of the PEO, m , was calculated by dividing the enthalpy
of fusion determined by DSC by the one corresponding to 100% crystalline PEO. Its
value is thus given by: m = Hm /Hm0 with Hm0 = 210 J/g.
Fig. 2. Scanning electron micrographs of cryo-fractured surface for cast-evaporated
nanocomposite electrolyte films reinforced with 6 wt% of sisal MF.
effects, which increase the apparent cross-section of whiskers [16].
For the composites reinforced with MF, the filler is observed with
more difficulty, which may be associated with the more chaotic
aspect of the surface.
3.1.2. Composite electrolytes
The morphology of the nanocomposite polymer electrolytes
reinforced with whiskers and MF was also characterized by SEM
observations. As nanocomposite electrolytes based on whiskers
were obtained by adding the salt after the composite film elaboration, the morphology of whiskers based composite electrolytes is
similar to that observed without salt.
For MF-based composite electrolytes, the salt was added before
the preparation of composite films and may have an influence on
the MF dispersion. However, as shown by the SEM micrography
(Fig. 2) of the nanocomposite electrolyte reinforced with 6 wt% MF
at a salt concentration O/Li equal to 12, the composite electrolytes
exhibited the same morphology than salt-free PEO–MF composites.
White dots associated with the MF entities are homogeneously
dispersed in the matrix with no apparent aggregates. The incorporation of salt has no effect on the MF dispersion which may be
controlled by PEO/MF interaction, dispersion process and solution
viscosity.
3.2. Thermal characterizations
3.2.1. Salt-free PEO nanocomposites
DSC measurements were performed on the PEO matrix and
related whiskers or MF reinforced composites. The DSC measurements were carried out 1 month after the film formation in order
to ensure an equilibrium state of crystallization to be reached and
to obtain reproducible results. All the characteristic temperatures
of the studied films are summarized in Table 2.
The melting temperature, Tm , is roughly constant and does not
dependence on the whiskers nature. The melting point of PEO
only decreases by adding MF. The crystallization temperature, Tc ,
decreases weakly with the incorporation of whiskers. Tc of composites is 1 or 3 ◦ C lower than that of the neat PEO matrix. The
shift of the crystallization process was supposed to result from the
affinity of the PEO with the reactive cellulose surface, restricting
locally the molecular mobility of polymer chains and their global
self-diffusion [18]. This mobility decrease may decrease the crystal
growth rate, thus the associated temperature. The largest effect on
Tc values was observed with the addition of sisal whiskers which
exhibit the highest specific area.
The crystallization temperature is notably affected by the presence of MF (Table 2). It may be ascribed to the entanglements of
these long fibers that can restrict the molecular mobility of polymeric chains.
The degree of crystallinity of the PEO matrix, m was calculated
using the ratio between the enthalpy of fusion determined by DSC
and the one corresponding to 100% crystalline PEO [19]. It was calculated per gram of PEO to take into account the presence of the
filler.
The incorporation of whiskers induces a small decrease of the
degree of crystallinity, m, at the equilibrium state. The cellulosic
nanoparticles most probably act as defects in the composite films,
limiting the organization of PEO chains and their ability to crystallize at the equilibrium state.
3.2.2.1. Melting temperature. The effect of the presence of both
salt and whiskers or MF on the PEO melting and crystallization
temperatures were evaluated by non-isothermal experiments. In
PEO–LiTFSI electrolytes, the salt induces two important effects:
(i) the introduction of ionic charge carriers needed for ionic conduction and (ii) the decrease of both the degree of crystallinity
and kinetic, through interactions between PEO chains and lithium
cations and bulky anions which interfere with the formation of PEO
regular lamellae.
The addition of 6 wt% nanocharges induces small modifications
of the PEO melting process in PEO–LiTFSI. Indeed, the composite
electrolytes exhibit the same melting temperature and nearly the
same melting enthalpy than PEO–LiTFSI (Table 3).
3.2.2.2. Glass transition temperature. The glass transition temperature, Tg , of the unfilled polymer electrolytes increases with the salt
concentration. This phenomenon is associated with the well know
stiffening effect of lithium salt on PEO chains. The addition of cellulosic fillers has no significant effect on Tg values, as previously
observed for PEO based composites [14,20]. In most of the studies
performed on inorganic composite electrolytes [21,10,22–26], the
effect of filler incorporation on Tg is small, which may be associated
with two antagonist effects, i.e. PEO amorphization which weakly
decrease the Tg value and specific filler/PEO interactions which may
increase the Tg value.
As the degree of crystallinity is not modified upon addition of
whiskers, the invariance of Tg may be related to the predominant
effect of lithium cation/PEO interactions compared to filler/PEO
interactions.
5190
Table 3
Thermal characteristics of PEO based nanocomposites reinforced with sisal (WS), ramie (WR) and cotton (WC) whiskers and sisal microfibrils (MF) obtained from DSC
curves: glass transition (Tg ), crystallization temperature (Tc ), melting temperature (Tm ), enthalpy of melting and enthalpy of crystallization related to the weight of PEO or
PEO-complex without charge.
Tg (◦ C)
Tm (◦ C)
Hm (J/gPEO–salt )
Tc (◦ C)
Hc (J/gPEO–salt )
PEO
PEO-O/Li = 20
PEO-O/Li = 12
−55
−51
−48
56
50
41
170
112
75
51
47
30
145
58
15
PEO-O/Li = 20 + 6 wt% WS
PEO-O/Li = 20 + 6 wt% WR
PEO-O/Li = 20 + 6 wt% WC
PEO-O/Li = 20 + 6 wt% MF
−52
−52
−50
−51
47
48
50
49
115
112
115
113
41
43
45
45
73
73
75
74
PEO-O/Li = 12 + 6 wt% WS
PEO-O/Li = 12 + 6 wt% WR
PEO-O/Li = 12 + 6 wt% WC
PEO-O/Li = 12 + 6 wt% MF
−48
−49
−46
−47
40
40
42
41
67
66
67
65
30
31
31
32
24
23
24
26
3.2.2.3. Crystallization process. The effect of the whiskers incorporation on the crystallization process, obtained during the cooling
ramp depends on the salt concentration (Fig. 3).
At low salt concentration, O/Li = 20, the crystallization temperature decreases with the whiskers incorporation. The decrease
observed depends on the whiskers nature, the highest effect being
obtained with sisal whiskers (Fig. 3). This behavior may be associated with the highest specific area of the sisal whiskers. For
O/Li = 12, no effect of the whiskers incorporation is observed. This
could be ascribed to the predominant influence of the salt on the
crystallization process.
For each salt concentration, the crystallization enthalpy of composite electrolytes, determined during the cooling ramp (Table 3),
Hc (J/gPEO–LITFSI ), is higher than that of unfilled electrolytes. While
the melting enthalpy of composite electrolytes, determined during
the heating ramp, Hm (J/gPEO–LITFSI ), is equal or lower than that of
unfilled electrolytes. These two informations indicate that cellulose
fillers induce an increase of the crystallization kinetic of the electrolytes as previously observed for salt-free PEO composites [14].
This behavior may be explained by the fact that the crystallization
rate is the product of the nucleation rate and the crystal growth.
Thus, the crystallization kinetic may increase despite a slowing
down of the crystal growth rate when more crystals are nucleated.
Indeed, Azizi Samir et al. [27] have shown a net decrease of the
PEO spherulites size and a large increase in the spherulites amount
with the addition of tunicin whiskers. Furthermore, the crystallization process is still incomplete, the Hc (J/gPEO–LITFSI ) is lower than
Hm (J/gPEO–LITFSI ), and the difference is larger for the filled and
unfilled electrolytes with a salt concentration O/Li = 12 than for the
salt-free samples.
Fig. 3. Normalized DSC thermograms showing the non-isothermal crystallization
at 10 ◦ C/min for PEO (), PEO–LiTFSI O/Li = 20 (+), PEO–LiTFSI O/Li = 20 + 6 wt%
sisal whiskers (), PEO–LiTFSI O/Li = 12 (×), and PEO–LiTFSI O/Li = 12 + 6 wt% sisal
whiskers (♦).
This thermal behavior was not obviously observed for nanocomposite PEO electrolytes based on inorganic charges for which
an amorphization of the composite electrolytes was generally
obtained [28,29] with a decrease of the crystallization kinetics. The
explanation given in the literature is that the inorganic nanocharges
prevent the local PEO organization, and thus reduce the crystallization kinetics. In cellulose whiskers/PEO composites, an increase
of the crystallization process was observed and it is ascribed to
an increase of the nucleation rate [14]. For cellulose whiskers/PEO
electrolytes, the same explanation can be promoted to explain the
enhancement of the crystallization process with the addition of
whiskers.
3.3. Degradation behavior
3.3.1. Nanocharges and salt-free PEO nanocomposites
The thermal stability of whiskers and MF was characterized
using thermogravimetric analysis under helium flow at 10 ◦ C/min
(Fig. 4). For all samples, the degradation process occurred in two
steps as already reported elsewhere for ramie whiskers [14] and
Kraft paper [30]. The first step around 250–300 ◦ C corresponds to
hemicellulose and glucosidic link depolymerization and the second steps, in the temperature range between 400 ◦ C and 430 ◦ C, is
attributed to the thermal degradation of ␣-cellulose [31].
For ramie whiskers, a higher thermal stability was obtained with
the onset degradation temperature, i.e weight loss, at 265 ◦ C. For
sisal whiskers and microfibrils, the onset degradation temperatures were very similar at 250 ◦ C and 244 ◦ C, respectively, whereas
for cotton whiskers the weight loss starts at a lower temperature,
around 229 ◦ C (Fig. 4). Roman and Winter [32] evaluated the influence of the acid charge density on the cellulose whiskers surface.
They reported a net decrease of the whiskers thermal stability when
increasing the acid charge density. In order to complete the data
Fig. 4. TGA curves measured under helium flow for (+) ramie, (♦) sisal, () cotton
whiskers, and () sisal MF.
Fig. 5. TGA curves measured under helium flow at 10 ◦ C/min for the salt-free unfilled
PEO matrix () and PEO + 6 wt% of sisal MF () ramie (♦) sisal (×), and cotton (−)
whiskers.
exhibited in the literature, i.e. 0.021 e/nm2 for ramie whiskers [15]
(e as electron charge, 1.6 × 10−19 C), the acid surface density was
determined for cotton and sisal whiskers by titration. Cotton and
sisal whiskers exhibit a surface charge density equal to 0.029 e/nm2
and 0.026 e/nm2 , respectively, thus higher than the value obtained
for ramie whiskers. The highest charge density of cotton whiskers
can explain the lower thermal stability observed. The thermal stability of whiskers under helium is clearly associated with the acid
charge density, and the lower the charge surface density is (Table 1),
the higher the thermal stability is.
The lower temperature value obtained for microfibrils compared to sisal whiskers was previously attributed to their higher
amorphous domain density and to the presence of pectins [13].
The study of the thermal degradation of whiskers and MF reinforced PEO films was carried out with a temperature ramp of
10 ◦ C/min under inert atmosphere, i.e. helium in order to avoid any
oxidizing character of the medium. The results are reported in Fig. 5
and Table 2.
Compared to the neat PEO, the presence of 6 wt% of cellulosic
nanoparticles increases the temperature of the main degradation,
associated with the polymer matrix. However, the degradation
occurs at lower temperature and was associated to the whiskers
degradation. The observed weight loss is low because of the low
whiskers content in the composite films.
The thermal stability of cellulose nanocrystal and MF reinforced
PEO–LiTFSI polymer electrolytes was investigated through thermogravimetric analysis (TGA) under helium flow at 10 ◦ C/min. Typical
TGA measurements are shown in Fig. 6, for unfilled PEO-O/Li = 12
and related composite electrolytes reinforced with 6 wt% filler.
Fig. 6. TGA measurements under helium flow at 10 ◦ C/min for unfilled () POEO/Li = 12, and related composites filled with 6 wt% (−) sisal MF, () sisal, (×) ramie,
and () cotton whiskers.
5191
Fig. 7. Evolution of the logarithm of the storage tensile modulus E’ vs temperature
at 1 Hz for neat PEO (), PEO + 6 wt% sisal (+), ramie () and cotton whiskers (×),
and sisal MF (−).
The thermal degradation phenomena of PEO and PEO–lithium
salt complexes were largely reported with complex mechanism. In
PEO and PEO–lithium complex, employing LiCl and LiI lithium salts,
the thermal degradation occurs in a single step at about 350 ◦ C.
Costa et al. [33] have remarked the complexity of thermal degradation mechanism invoking a strong interaction between the metal
ion and the oxygen atoms of the polymer.
The incorporation of the filler decreases the onset thermal
degradation for all samples. A degradation process with two steps
was clearly observed. The first step, around 300 ◦ C, as for salt-free
PEO composites, is associated with the whiskers degradation. The
lower temperature onset is observed for cotton based composite,
in accordance with the fact that cotton whiskers present the lowest thermal stability. The main thermal degradation, around 420 ◦ C
for filled electrolytes, associated with PEO complex degradation, is
obtained at higher temperature than the one observed for unfilled
PEO electrolytes, except for cotton whiskers based composite in
accordance to thus obtain for salt-free samples. The degradation
temperature increase induces by the filler incorporation is more
important than the one observed for salt-free samples.
3.4. Dynamic mechanical analysis.
The mechanical behavior of nanocomposite polymers and electrolytes was evaluated in the linear range using DMA under
isochronal condition at 1 Hz. At low temperature, T = −70 ◦ C, the
storage modulus was normalized at 7.7 GPa for all samples, which
corresponds to the observed average value regardless the composition. This normalization leads to minimize the influence of the
inaccurate dimensions of samples on storage modulus.
3.4.1. Salt-free PEO composites
The mechanical behavior for salt-free PEO composites was
investigated to compare the effect of the various cellulosic fillers
(Fig. 7). Above the main relaxation process associated with the
glass–rubber transition, around −55 ◦ C, the composite filled with
sisal whiskers exhibits a higher storage tensile modulus. The
increase of the stiffness of the composite is a mark of the reinforcement effect induced by the presence of the sisal whiskers. For
other samples, such strengthening was less important.
Above the melting point, 65 ◦ C, the storage tensile modulus
fall down for all samples. The matrix melts and a plateau appears
thanks to the cellulose filler. For composites reinforced with MF, the
temperature for which the modulus drops is shifted towards about
85 ◦ C. A similar behavior was already reported by Ogata et al. [34]
and was ascribed to high filler/matrix interactions or MF entanglements. PEO composites reinforced with MF display a storage tensile
modulus around 16 MPa associated with an entanglement effect of
MF within the matrix.
5192
Table 4
High temperature (T = 80 ◦ C) tensile modulus: comparison between experimental
(E’exp ) and predicted (E’pre ) data for whiskers reinforced PEO nanocomposite films.
Sample
Cotton
whiskers 6 wt%
E’exp (MPa)
E’pre (MPa)
7
–
Ramie
whiskers 6 wt%
12
4
Sisal whiskers
6 wt%
160
97
At high temperature, the highest reinforcement was obtained
for sisal whiskers with 120 MPa up to 160 ◦ C. For other samples,
the reinforcement is less important and disappears at lower temperature. This behavior could be associated with the formation of a
whiskers network which strongly depends on the aspect ratio and
the modulus of the whiskers percolating network.
The modulus of this continuous network reinforced composite
can be well predicted from the adaptation of the percolation concept to the classical series-parallel model [35]. In this model and
at sufficiently high temperature, i.e. when the storage modulus of
the matrix is much lower than the one of the percolating network,
the following equation [36] was derived for the predicted elastic
modulus, E’pre , of the composite:
Epre
= ER
(1)
With:
=0
− b
R
Rc
= R
1 − Rc
for R < Rc
for R < Rc
(2)
where and E’R are the volume fraction and the elastic modulus
of the rigid percolating network, respectively; R , Rc and b correspond to the volume fraction of filler, critical volume fraction of
filler at the percolation threshold and corresponding critical exponent, respectively. For a 3D network, b = 0.4 [37] and Rc = 2.5 vol%,
1.6 vol%, 5.4 vol% were determined from the aspect ratio of ramie,
sisal and cotton whiskers, respectively. E’R can be experimentally
determined from a tensile test performed on a film prepared by
water evaporation of a suspension of cellulose whiskers. The tensile
modulus of a ramie whiskers film, E’R , was determined elsewhere
[15] and a value of 0.35 GPa was reported. For sisal whiskers, the
modulus of the percolating network, E’R was experimentally determined and found around 8.5 GPa. The E’R value of a film made of sisal
whiskers is one of the highest values compared to ramie whiskers
or tunicin whiskers, between 5 GPa and 15 GPa [35,38].
The predicted storage modulus values, E’pre , are reported in
Table 4. It was not determined for the nanocomposite film reinforced with 6 wt% cotton whiskers, because this filler content
is slightly lower than the theoretical value for the percolation threshold. However, experimentally, a stabilization of the
storage modulus was observed with 6 wt% of cotton whiskers
(Rc = 4.9 vol%), thus very close to the theoretical percolation
threshold. This difference may be associated with the model developed, which neglects the effect of whiskers reinforcement below
the percolation threshold.
In accordance with the predicted elastic modulus, the sisal
whiskers based composite exhibited the highest experimental
storage modulus at high temperature. The lowest mechanical properties obtained with cotton whiskers at high temperature can be
associated with the too low amount of whiskers, i.e. below the
percolating threshold.
For the evaluation of composite electrolytes, we have chosen
PEO films reinforced with 6 wt% sisal whiskers and MF. Indeed,
PEO–6 wt% sisal whiskers exhibits the highest reinforcing effect
for salt-free PEO nanocomposites compared to ramie and cotton
whiskers, in agreement with the highest aspect ratio and storage
modulus of sisal whiskers.
Fig. 8. Storage tensile modulus E’ vs temperature at 1 Hz for PEO (-), PEO-O/Li = 20
(+), PEO-O/Li = 12 (×), PEO-O/Li = 12 + 6 wt% MF (), and PEO-O/Li = 12 + 6 wt%
whiskers sisal ().
3.4.2. Nanocomposite polymer electrolytes
The addition of salt induces a net decrease of the tensile modulus above the glass–rubber transition of PEO, associated with an
increase of the PEO amorphous phase content (Fig. 8). Above −55 ◦ C,
PEO-O/Li = 12 exhibits a lower storage tensile modulus than the
PEO-O/Li = 20 complex, which is associated with the less crystalline
structure of this electrolyte.
In the rubbery plateau zone, the mechanical stifness of filled
electrolytes is due to both the cellulosic filler and PEO crystalline
domains. The addition of the filler induces a decrease of the degree
of crystallinity at the equilibrium state, with 75 J/gPEO–salt and near
66 J/gPEO–salt for PEO–LiTFSI O/Li = 12 and filled PEO–LiTFSI O/Li = 12,
respectively (Table 4).
This lower crystallinity induces a decrease of the storage modulus. This decrease can be at least compensated by the reinforcing
effect of the percolating whiskers network. As the reinforcing effect
of MF is lower than the one of sisal whiskers, a decrease of the
rubbery storage modulus was observed for the MF filled electrolyte.
For pure PEO and low salt concentration the storage tensile modulus drops irreversibly when the temperature reaches the melting
point.
For a higher salt concentration (O/Li = 12), a decrease of the
storage tensile modulus was observed at the melting temperature,
but it remains constant around 2 MPa for high temperatures up to
145 ◦ C. This behavior is related to lithium cation/PEO interactions
which involve ionic cross-linking between PEO chains. For lower
salt concentrations, O/Li = 20, the ionic cross-links density is most
probably not sufficient to permit transient network behavior above
the melting point.
For the PEO-O/Li = 12 complex, a spectacular increase of the
tensile modulus is observed with the incorporation of 6 wt% of
sisal whiskers (Fig. 8). The tensile modulus of PEO + 6 wt% of sisal
whiskers remains constant up to 140 ◦ C with a value higher than
120 MPa. This reinforcement is attributed to the formation of a rigid
percolation whiskers network as reported for salt-free PEO based
composites. It is an indication that salt addition in the composite
doesnot modify the cohesion of the whiskers network.
The same behavior was reported by Azizi Samir et al. [18]
for nanocomposite electrolytes PEO–LiTFSI reinforced with tunicin
whiskers. For these nanocomposites, and above the melting point
of PEO, the thermal stabilization of the storage tensile modulus was
observed regardless the salt content, with a modulus value that was
found to decrease as the salt concentration increased.
The electrolyte film PEO + O/Li = 12 + 6 wt% sisal whiskers
exhibits lower mechanical properties than the salt-free sample,
i.e. 120 MPa compared to 160 MPa at the same temperature. This
decrease may be explained by a decrease of the effective amount
of cellulosic filler in the composite polymer electrolyte after salt
Fig. 9. Arrhenius plot of the ionic conductivity for polymer electrolyte PEO–LiTFSI
complex with O/Li = 20 for (♦) unfilled electrolyte and filled electrolyte with 6 wt%
(×) sisal, () ramie, (+) cotton whiskers, and () sisal MF.
addition. Indeed, as it was noted in Section 2, whiskers reinforced
nanocomposite PEO films were first prepared and lithium salt was
added later by film swelling. Therefore, the effective amount of sisal
whiskers, calculated as a ratio between the weight of whiskers and
that of the composite becomes 4.9 wt% instead of 6 wt% for salt-free
samples. It is, then, possible to compare this experimental modulus
value with the predicted one using this real whiskers content. The
predicted data were calculated as previously done for salt-free PEO
composites. For the calculation, the density of crystalline cellulose
was taken as 1.5 g cm3 and the one of the polymer electrolyte matrix
and LiTFSI were taken as 1.2 g cm3 and 2.33 g cm3 , respectively. The
density of the polymer electrolyte O/Li = 12 was therefore taken as
1.45 g cm3 . The predicted modulus was found to be 72 MPa.
The values of the ratio E’exp /E’pre for both PEO–6 wt% sisal
whiskers and PEO-O/Li = 12 + 6 wt% sisal whiskers are both equal
to 1.7. This observation confirms that the difference observed can
be ascribed to the decrease of the effective sisal whiskers content
during the preparation of the electrolyte.
With a salt concentration O/Li = 12, the storage tensile modulus
obtained with 6 wt% sisal MF is close to 15 MPa (Fig. 8). This value is
similar to the one obtained for salt-free PEO–6 wt% MF composite,
i.e. 16 MPa. The presence of the salt does not modify the mechanical
properties of the PEO–MF-based composites. Indeed, the reinforcing effect is governed by both the tensile strength of cellulose
microfibrils via entanglements and PEO/microfibrils interaction.
3.5. Ionic conductivity
3.5.1. Whiskers reinforced nanocomposites
In order to avoid proton conduction due to the presence of acidic
charge at the whisker surface, the acidic functions were neutralized
by LiOH.
The thermal dependence of the ionic conductivity for composite
PEO polymer electrolytes with O/Li = 20 is reported in Fig. 9. The
temperature sweep test was carried out from 25 ◦ C to 75 ◦ C and
then from 75 ◦ C to 20 ◦ C. Fig. 9 reports the ionic conductivity values
measured when decreasing the temperature from 75 ◦ C to 20 ◦ C.
The ionic conductivities for PEO/LiTFSI complexes with O/Li = 20
are higher than 10−4 S cm−1 for T > 45 ◦ C.
The highest conductivity was obtained with the addition of
sisal whiskers, with a conductivity value of 3.1.10−4 S cm−1 at 60 ◦ C
compared to 4.10−4 S cm−1 for PEO–LiTFSI O/Li = 12. Thus, the conductivity decrease is very weak especially at high temperature.
The conductivity value obtained for films reinforced with cotton whiskers is the lowest one. DSC investigations (Table 4) have
shown that the glass transition temperature was not affected by
the addition of whiskers, thus a macroscopic restriction of the
5193
PEO mobility cannot explain the observed behavior. A decrease
of the ionic conductivity with the addition of tunicin whiskers
in PEO based electrolytes was also observed by Azizi Samir et al.
[18] Using NMR investigation, this phenomenon was ascribed to (i)
the reduction of the relaxation time of polymer chain and (ii) the
decrease of diffusion coefficient of both anion and cation. These
two results seem to indicate that the decrease in ionic conduction
may be associated with the restriction of electrolyte mobility at the
whiskers/electrolyte interface and not significantly in the bulk (i.e.
invariance of Tg ). Such behavior was observed for PEO–TiO2 composites using quasielastic neutron scattering [39]. The existence of
an immobilized layer of polyether matrix around the TiO2 particles was suggested, whereas the bulk electrolyte properties were
unaffected by the presence of the filler. However, the decrease
of conductivity is not directly associated with the whiskers specific surface, as the lowest conductivities were obtained for cotton
whiskers which display the lowest specific area.
Below the complex melting temperature, the difference
between the conductivity of the unfilled and filled electrolytes
becomes larger. The net decrease of conductivity observed at lower
temperatures is due to the presence of the crystalline phase. The
difference observed between filled and unfilled samples may be
due to an increase of the crystallization kinetic in the presence of
whiskers, in agreement with DSC measurements, i.e. higher crystallization kinetics observed during the DSC cooling ramp (Table 3).
This behavior is not commonly observed in composite electrolytes.
Generally, an improvement of the ionic conductivity is obtained at
room temperature [40] and a quasi invariance [41] is observed at
higher temperatures. This difference is associated with the crystallization process, which is slowed down with the addition of
inorganic nanoparticles and accelerated with whiskers.
3.5.2. Microfibrils reinforced nanocomposites
The addition of sisal microfibrils induces a decrease of the ionic
conduction, by a factor close to 3 at 60 ◦ C compared to PEO–LiTFSI
O/Li = 20. This larger effect compared to sisal whiskers may be
ascribed to the addition of long and entangled fibrils, which may
restrict more notably the mobility of PEO chains.
4. Conclusions
The effect of cellulose microfibrils and whiskers as a reinforcing phase in polymer electrolytes was investigated. The salt
used was LiTFSI. For a given salt concentration, Tg , Tm and Hm
were found to be independent of the whiskers content. Nevertheless, the incorporation of cellulose whiskers has an influence on
Tc and Hc for low salt concentration, O/Li = 20. The addition of
whiskers and MF in polymer electrolytes leads to high performance
nanocomposites with a high increase in the storage modulus at high
temperature, compared to the unfilled PEO–LiTFSI sample. Ionic
conductivity measurements have shown that, for a given temperature, the presence of whiskers or MF induces a weak decrease of
polymer electrolyte conductivity, which may be due to the existence of interactions between cellulose and PEO or lithium salt. The
small decrease in ionic conduction is largely compensated by a high
reinforcing effect, especially at high temperature, obtained with
the use of natural, biodegradable, low density and largely available
nanofibers.
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Cellulose
DOI 10.1007/s10570-011-9543-x
Poly(oxyethylene) and ramie whiskers based
nanocomposites: influence of processing: extrusion
and casting/evaporation
Fannie Alloin • Alessandra D’Aprea
Alain Dufresne • Nadia El Kissi •
Frédéric Bossard
•
Received: 21 May 2010 / Accepted: 5 April 2011
Ó Springer Science+Business Media B.V. 2011
Abstract Polymer nanocomposites were prepared
from poly(oxyethylene) PEO as the matrix and high
aspect ratio cellulose whiskers as the reinforcing phase.
Nanocomposite films were obtained either by extrusion or by casting/evaporation process. Resulting films
were characterized using microscopies, differential
scanning calorimetry, thermogravimetry and mechanical and rheological analyses. A thermal stabilization
of the modulus of the cast/evaporated nanocomposite
films for temperatures higher than the PEO melting
temperature was reported. This behavior was ascribed
to the formation of a rigid cellulosic network within the
matrix. The rheological characterization showed that
nanocomposite films have the typical behavior of solid
materials. For extruded films, the reinforcing effect of
whiskers is dramatically reduced, suggesting the
absence of a strong mechanical network or at least,
A. D’Aprea N. E. Kissi F. Bossard
Laboratoire de Rhéologie, UMR 5520, Grenoble-INPCNRS-UJF, BP 53, 38041 Grenoble Cedex 9, France
F. Alloin (&) A. D’Aprea
LEPMI, Laboratoire d’Electrochimie et de Physicochimie
des Matériaux et des Interfaces, UMR 5279, CNRSGrenoble INP-Université de Savoie - Université Joseph
Fourier, BP 75, 38402 Grenoble Cedex 9, France
e-mail: [email protected]
A. D’Aprea A. Dufresne
The International School of Paper, Print Media
and Biomaterials (Pagora), Grenoble-INP, BP 65,
38402 Saint Martin d’Hères Cedex, France
the presence of a weak whiskers percolating network.
Rheological, mechanical and microscopy studies were
involved in order to explain this behavior.
Keywords Polymer matrix composites Cellulose
whiskers Thermal properties Dynamic mechanical
analysis Rheology Extrusion
Introduction
During the last decade, natural fibers reinforced
thermoplastic polymers have attracted the attention of
both the academic and industrial world for applications in transport and construction. This considerable
interest in the possibility of replacing conventional
fibers, like glass fibers, is ascribed to well-known
advantages such as renewable nature, low cost and
density. However, one of the main drawbacks of
lignocellulosic fibers is the big variation of properties
inherent to any natural products. Their properties are
related to climatic conditions, maturity, and type of
soil. Disturbances during plant growth also affect the
plant structure and are responsible for the enormous
disparity of mechanical plant fiber properties. One of
the basic idea to achieve further improved fiber and
composite is to eliminate the macroscopic flaws by
destructuring the natural grown fibers, and separating
the almost defect free highly crystalline fibrils. When
combining this process with an acidic treatment, high
123
Cellulose
specific surface area rod-like nanoparticles of monocrystalline cellulosic fragments can be obtained. The
state of dispersion of these nanoparticles, called
cellulose whiskers, in a polymeric matrix has an
important impact on the final properties of composites and is strongly dependent on the processing
technique and conditions.
In a previous study (Azizi Samir et al. 2004a),
poly(oxyethylene) (PEO) films have been reinforced
with tunicin whiskers. PEO being a hydrosoluble
polymer and cellulose whiskers being obtained as
aqueous suspensions, the processing of nanocomposite films was easily carried out by mixing the two
constituents in water and casting the resultant
dispersion. The important aspect ratio of tunicin
whiskers make these cellulose nanoparticles derived
from tunicate, a sea animal, a good candidate for
modeling rheological and reinforcement behaviors
and it was extensively studied in the literature (Favier
et al. 1995). However, this source of cellulose is not
suitable for technical applications due to its poor
availability. Such rod-like nanoparticles with various
aspect ratios can also be extracted from other
renewable resources as plant fibers (Azizi Samir
et al. 2005a). In the present study, cellulose whiskers
were extracted from ramie fibers.
Very few studies have been reported concerning
the processing of cellulose whiskers reinforced
nanocomposites by extrusion methods. An attempt
to prepare nanocomposites based on cellulose whiskers obtained from microcrystalline cellulose and
poly(lactic acid), PLA, by melt extrusion technique
was recently reported (Oksman et al. 2006; Bondeson
and Oksman 2007). The suspension of nanocrystals
was pumped into the polymer melt during the
extrusion process. Organic acid chlorides-grafted
cellulose whiskers were also extruded with low
density polyethylene (De Menezes et al. 2009).
In the present study, we investigate the preparation
and the properties of nanocomposite films obtained
from PEO as the matrix and cellulose whiskers
extracted from ramie as the reinforcing phase. Both
casting/evaporation, largely employed in research,
and extrusion processing method, a more industrial
technique, have been used. Thermogravimetric analysis (TGA) and differential scanning calorimetry
(DSC) measurements have been used to investigate
the thermal characteristics and degradation of the
ensuing nanocomposites. The viscoelastic behavior of
123
these materials was investigated in both the molten
and solid states as a function of the whiskers content
and processing technique by rheological and mechanical methods. In the first part of the paper, the effect
of the ramie whiskers content on the properties for
cast/evaporated nanocomposite films was investigated. In the second part, the impact of the processing
technique was examined, restricting the study to a
film reinforced with 6 wt% ramie whiskers.
Experimental methods
Nanocomposite films
Materials
Poly(oxyethylene), PEO, with a high molecular
weight (Mw = 5 9 106 g mol-1) was purchased as
a white powder from Aldrich and used as received.
Cellulose whiskers were prepared from ramie fibers
as described in details elsewhere (Habibi and Dufresne
2008). Briefly, ramie fibers were cut into small pieces
and treated with 2 wt% NaOH at 80 °C for 2 h to
remove residual additives. The purified ramie fibers
were submitted to acid hydrolysis with a 65 wt%
H2SO4 solution at 55 °C for 30 min and under continuous stirring. The suspension was washed with water by
centrifugation and dialyzed to neutrality against deionized water. The obtained suspension was homogenized
using an Ultra Turrax T25 homogenizer for 5 min at
13,500 rpm and then filtered in sintered glass No. 1 to
remove unhydrolyzed fibers. The suspension was
concentrated to constitute the stock suspension. This
treatment leads to aqueous suspensions of high aspect
ratio rod-like nanocrystals, characterized by an average
dimension of the cross section d of 6–8 nm and a length
L ranging between 150 and 250 nm, measured by TEM
(Habibi et al. 2007; Habibi and Dufresne 2008). The
average aspect ratio L/d and the specific surface area of
these whiskers were estimated to be close to 28 ± 12
and 380 ± 38 m2 g-1, respectively, taking 1.5 g cm-3
for the density of cellulose.
Processing of nanocomposite films
Cast/evaporated films The desired amount of ramie
whiskers aqueous suspension was added to the PEO,
previously dispersed in a few milliliter of methanol
Cellulose
for a better dissolution. The cellulose whiskers
content was varied from 0 to 30 wt% (dry basis).
The resulting suspension was protected against light
by an aluminum foil and was weakly stirred for
4 days at room temperature. The suspension was then
degassed under vacuum in order to remove the
remaining air, cast into Teflon plates and dried under
argon at 40 °C for 3 weeks. The films were then
progressively dried under vacuum for a week with a
temperature increase step of 5 °C per day from 40 to
75 °C and finally stored in glove box. The thickness
of final films was about 200–300 lm.
negative coloration which emphasizes cellulose. The
accelerating voltage was 80 kV.
A ZEISS polarizing optical microscope was used
to observe and follow the growth of the PEO
spherulites. The microscope was connected to a
LINKAM 20 to control the temperature. Samples
were melted at 100 °C for few minutes and then
cooled to the crystallization temperature. Average
radius values of spherulites were measured assuming
a circular shape.
Extruded films Poly(oxyethylene) matrix and ramie
whiskers reinforced PEO nanocomposite films were
prepared by extrusion. First, the PEO solution and
cellulose whiskers/PEO suspensions with 6 wt%
ramie whiskers were prepared similarly to the
previous description. The suspension was degassed
under vacuum and the water was removed by freezedrying. The ensuing powder was introduced in the
mixing chamber of a twin-screw DSM Micro 15
compounder and allowed to melt under nitrogen flow
at 180 °C. The mixing speed was set at 25 rpm for
10 min. Extrusion was carried out with a slit die of
0.6 mm in gap and 1 cm in length. The extruded
films were then cooled and calendered. They were
homogeneous, smooth, and bubble-free. The film
thickness ranged between 400 and 500 lm. The films
were dried for 4 days at 75 °C under vacuum and
then stored in glove box.
Differential scanning calorimetry tests were performed using a TA Instrument DSC, DSC2920 CE.
Standard modes were performed. Samples of 10 mg
were sealed in aluminum pans and placed in the DSC
cell in glove box. Each sample was heated from -100
to 100 °C at a temperature ramp of 10 °C min-1 and
kept at this temperature for 5 min to insure the
thermal equilibrium. Then, it was cooled down to
0 °C at a temperature ramp of 10 °C min-1. The
melting temperature, Tm, and the crystallization
temperature, Tc, were taken at the onset of the
melting and crystallization peaks, respectively.
Characterizations
Microscopies
Scanning electron microscopy (SEM) was used to
investigate the morphology of the nanocomposite
films using a LEO S440 SEM instrument. The
samples were frozen under liquid nitrogen, fractured,
mounted, coated with graphite and observed using an
accelerating voltage of 10 kV.
Transmission electron microscopy (TEM) observations were made with a Philips CM200 electron
microscope. A droplet of a dilute suspension of redispersed extruded nanocomposite films with 6 wt%
of cellulose whiskers was deposited and dried on a
carbon coated grid with one droplet (*6 ll) of
uranyl acetate solution (2 wt%) to carry out the
Differential scanning calorimetry
Thermogravimetric analysis
Thermogravimetric analysis measurements were carried out with a Netzsch STA409 thermal analyzer.
Around 10 mg of the sample were heated from room
temperature up to 550 °C at 10 °C min-1 under
either air or nitrogen flow. The results allow following the weight loss as a function of the temperature.
The degradation temperature was associated with the
beginning of the weight loss.
Dynamic mechanical analysis
Dynamic mechanical analysis (DMA) measurements
were carried out with a spectrometer DMA Q800
from TA Instrument working in the tensile mode. The
strain amplitude was fixed at 0.05%, well below the
limit of the linear viscoelastic regime. The samples
were thin rectangular strips with dimensions of about
20 9 7 9 0.2 mm3 for cast/evaporated films and
20 9 7 9 0.4 mm3 for extruded films. Measurement
of the storage tensile modulus, E0 , was performed in
isochronal condition (1 Hz), and the temperature was
123
Cellulose
varied between -100 and 150 °C using a temperature
ramp of 2 °C min-1.
Rheometry measurements
Rheometrical data were collected with a rotating
ARG2 rheometer from TA Instrument operating
under controlled strain conditions. It was equipped
with an oven and the analysis was carried out up to
180 °C under nitrogen flow. The storage shear
modulus G0 and the loss shear modulus G00 of the
samples were measured in the linear viscoelastic
regime using a parallel plates geometry of 25 mm in
diameter, with a gap of 500 lm. The dynamic
viscosity was calculated from G0 and G00 moduli.
Creep measurements for cast/evaporated and
extruded nanocomposites were performed on the
sample at 90 °C and at the same stress, corresponding
to a torque of 5 lNm.
Results and discussions
Cast/evaporated films
Morphology
The PEO films reinforced with ramie whiskers were
characterized by scanning electron microscopy (SEM).
Figure 1 shows the cryofractured surface for the cast/
evaporated unfilled PEO matrix and nanocomposite
film filled with 6 wt% ramie whiskers. The surface of
the cast/evaporated PEO matrix (Fig. 1a) appears
homogeneous without voids. The cryofractured
surface of the cast/evaporated nanocomposite film is
more chaotic than the matrix and shows a homogeneous dispersion of white dots (Fig. 1b). The cross
sections of these dots do not correspond to the one of
isolated whiskers, since their dimensions are far higher
than those of the whiskers. Indeed, it was shown that
the white dots result from electrical charge effects that
increase the apparent cross section of whiskers (Anglés
and Dufresne 2000).
Thermal characterization
The melting process was studied during the first
heating scan. For unfilled specimens, measurements
were performed for PEO powder and for the neat
cast/evaporated PEO matrix. No significant difference was reported between both unfilled materials as
can be seen in Table 1.
Melting temperature The melting temperature, Tm,
remains roughly constant for low whisker content, up
to 6 wt% (Table 1). At high filler content, i.e. 20 and
30 wt%, a weak decrease in Tm is observed. This
behavior is in good agreement with that reported by
(Azizi Samir et al. 2004a) for tunicin whiskers and
(Guo and Liang 1999) for wheat straw cellulose
whiskers. The decrease of the melting temperature
may be associated to the decrease of the spherulite
size. Indeed, a large decrease of the spherulite size
was observed with the incorporation of tunicate
whiskers in PEO matrix (Azizi Samir et al. 2005b).
Degree of crystallinity The degree of crystallinity
of the PEO matrix, vm, was calculated using the ratio
Fig. 1 Scanning electron micrographs of the cryofractured surface for the cast/evaporated a PEO film and b PEO film reinforced
with 6 wt% of ramie whiskers
123
Cellulose
Table 1 Thermal characteristics of cast/evaporated PEObased nanocomposites reinforced with ramie whiskers obtained
from DSC and ATG curves: crystallization temperature (Tc),
Tc (°C)
Samples
melting temperature (Tm), degree of crystallinity expressed as a
function of the matrix weight measured during the melting
(vm), degradation temperature (Tonset), activation energies
Tm (°C)
vam
Tonset (10 °C min-1)Air
Tonset (10 °C min-1)He
Ea (kJ mol-1)
–
Ramie W.
–
–
–
270
265
PEO powder
51
57
0.82
–
350
380
PEO
51
56
0.81
205
–
–
PEO ? 3 wt% WR
50
56
0.82
194
352
205
PEO ? 6 wt% WR
50
57
0.80
191
357
210
PEO ? 10 wt% WR
47
56
0.80
188
345
190
PEO ? 20 wt% WR
46
53
0.76
187
320
170
PEO ? 30 wt% WR
45
53
0.75
185
310
180
a
mPEO
vm ¼ DHDH
where DH0m = 210 J g-1 is the heat of melting for 100% crystalline PEO
0
m
between the melting enthalpy determined by DSC
and the one corresponding to 100% crystalline PEO.
It was calculated per gram of PEO to take into
account the presence of the filler. The effect of the
whiskers content on the degree of crystallinity of
PEO is weak and it is only observed for highly filled
specimens for which a slight decrease of the degree
of crystallinity is observed upon ramie whiskers
addition.
Degradation behavior
The thermal stability of ramie whiskers reinforced
PEO nanocomposites was characterized using thermogravimetric analysis. In these experiments, the
weight loss of cast/evaporated nanocomposite films
was plotted in Fig. 2 as a function of temperature
under air flow with a temperature ramp of
10 °C min-1.
PEO matrix and whiskers degradations
The onset degradation temperature, i.e. the temperature associated with the beginning of the weight
loss, was found to be 205 °C for the unfilled PEO
matrix (Table 1). The PEO degradation under oxygen
atmosphere induces the formation of a large number
of volatile products involving several mechanisms
(Costa et al. 1992).
The thermo-oxidative degradation of pure ramie
whiskers was also investigated and the degradation
temperature was found to be 270 °C (Table 1;
Fig. 2). The degradation of ramie whiskers occurs
Fig. 2 Weight loss of unfilled PEO matrix (open circle), ramie
whiskers (filled square) and related nanocomposite films
reinforced with 3 (filled triangle), 6 (inverted triangle), 10
(diamond), 20 (open triangle) and 30 wt% (right-pointing
triangle) of ramie whiskers versus temperature. Measurements
were performed under continuous flow of dry air
following two steps. The first one at low temperature
may correspond to the cellulose depolymerization
induced by the glucosidic bond scission which
involves hemicellulose formation, and the second
step, at about 400 °C, is most probably associated
with the thermal degradation of the a-cellulose, by
similarity with the degradation process reported for
sisal fibers (Alvarez et al. 2004).
Composite degradation under air flow
The weight loss observed for filled PEO films starts at
lower temperature compared to neat PEO, from 194
to 185 °C for the composites and at 205 °C for neat
123
Cellulose
PEO (Table 1; Fig. 2). A significant decrease of the
onset degradation temperature of the filled matrix is
observed even at low whisker content. It is roughly
25–30 °C lower than the onset degradation temperature of the neat polymer. Furthermore, the weight
loss kinetic is notably enhanced in the presence of
ramie whiskers (Fig. 2). Indeed, the quasi-total
degradation of the filled samples is observed at
300 °C, whereas it is achieved at 400 °C in the case
of neat PEO.
These results indicate a large effect of the presence
of whiskers on the PEO thermal stability even if the
whisker degradation occurs at high temperature.
(Chauvin et al. 2006) have shown that PEO is very
sensitive to acid medium and a strong decrease of its
molecular weight was observed. Furthermore, a net
decrease of the thermal stability of oligoether sulfate
(Chauvin et al. 2005) was observed in the presence of
a small amount of water, indicating the sensibility of
PEO degradation to the presence of both water and
sulfate acidic function. Thus, the sulfate ester groups
present at the whisker surface and resulting from the
acid hydrolysis treatment with sulfuric acid can have
a strong influence on the filled PEO degradation.
At low acid concentration, the sulfate ester groups
induce a significant decrease of the cellulose thermal
stability under air flow (Roman and Winter 2004;
Kim et al. 2001; Julien et al. 1993; Parks 1971; Tang
and Neill 1964). During the first degradation step
starting at 150 °C, the desulfation and dehydration of
the cellulose occur (Guo and Liang 1999). The latter
effect increases the water content in the medium and
catalyzes the degradation reaction of cellulose (Scheirs et al. 2001).
which may enhance the PEO degradation by hydrolysis. Indeed, the ether oxygen could provide
hydroxyl groups using acid catalysis, and then the
dehydratation of the PEO occurs (Grassie and
Mendoza 1985). The degradation of PEO is enhanced
by its oxidation induced by the presence of O2.
The activation energy, Ea, of this degradation
process can be determined using the Broido’s method
(Broido 1969) as follows:
1
Ea 1
ln ln
þC
ð1Þ
¼
y
R T
where y is the fraction of degraded product at time t,
R, T and C are the gas constant, the temperature and a
temperature independent term, respectively. Ea/R is
given by the slope of the plot of ln (ln 1/y) as a
function of 1/T. The activation energy obtained for
filled PEO is much lower than that obtained for neat
PEO (Table 1). This result suggests that the degradation reaction is initiated by the whiskers degradation. The PEO degradation is then catalyzed by the
sulfuric acid and water present in the medium.
In order to reduce the influence of sulfate acid
groups brought by the cellulose whiskers on the
thermal degradation of filled PEO, the whisker
suspension was neutralized by NaOH solution. A
nanocomposite film reinforced with 6 wt% of ramie
whiskers was prepared using these neutralized whiskers. TGA results are reported in Fig. 3. The onset
degradation temperature is similar for nanocomposite
Effect of acidic surface density
In order to evaluate the influence of sulfate acid
groups, present at the whisker surface, on the thermal
degradation of PEO based composites, the surface
density of sulfate groups of the ramie whiskers was
determined by titration using NaOH solution. It was
equal to 0.022 e/nm2 ± 0.001, taking 380 m2g-1 as
the specific surface area of the ramie whiskers. This
value is low but appears to be sufficient to induce a
decrease of the thermal stability of the cellulosic
nanoparticles (Roman and Winter 2004) without a
significant weight loss. This degradation induces the
formation of sulfate acid and water in the medium,
123
Fig. 3 Weight loss of the unfilled PEO matrix (filled triangle),
and related composites reinforced with 6 wt% of untreated
ramie whiskers (filled square) and neutralized ramie whiskers
(open circle) versus temperature. Measurements were performed under continuous dry air flow
Cellulose
Composite degradation under helium flow
was reported for a poly(propylene) matrix filled with
sisal microfibrils (Panaitescu et al. 2008).
Mechanical behavior
The temperature dependence of the storage tensile
modulus E0 for the unfilled cast/evaporated PEO
matrix and related composites is shown in Fig. 4. In
order to minimize the effect of the sample dimension
uncertainty on the accuracy of the measurement, the
glassy modulus, E0 at -100 °C, was normalized at
8.5 GPa for all the samples. It corresponds to the
observed average value regardless the composition of
the film. The unfilled PEO displays a typical behavior
of semi-crystalline polymer.
Normelized storage modulus (MPa)
films reinforced with neutralized or untreated whiskers (Fig. 3). However, the degradation kinetic
observed with neutralized whiskers based composite
is lower than that obtained with untreated whiskers
based composite.
The degradation activation energy of the nanocomposite film reinforced with 6 wt% neutralized
whiskers is about 217 kJ mol-1 which is notably
higher than the value obtained for the nanocomposite
film reinforced with untreated whiskers (190 kJ/
mol-1) but lower than the value obtained for neat
PEO (380 kJ/mol-1). The influence of acidic functions at the surface of the whiskers towards the
degradation of the PEO matrix is clearly evidenced;
however the addition of neutral whiskers has also a
negative impact on the PEO matrix thermal stability.
During the whiskers neutralization, the acid function
is transformed in a sodium salt. It has been reported
that alkaline salts have an influence on the degradation of PEO (Costa et al. 1992). The strong interaction between the cation and the ether oxygen involves
the weakening of the C–O bond and favors its
scission under air flow. In filled PEO, such interactions may occur in addition to hydrogen bonding
between cellulose OH groups and PEO. These
interactions may explain the decrease in thermal
stability observed in comparison to the neat PEO.
10
4
10
3
10
2
10
1
10
0
-100
-50
0
50
100
150
100
150
Normalized storage tensil modulus E' (MPa)
Temperature (°C)
The study of the thermal degradation of ramie
whiskers reinforced PEO films was carried out under
inert atmosphere, i.e. helium, to avoid any oxidizing
character of the medium. Results are reported in
Table 1. The onset degradation temperature of PEO
films, initially at about 205 °C in the presence of air
is shifted to 350 °C under helium, showing the
dominating effect of oxygen in the PEO degradation
mechanism (Cameron et al. 1989), whereas no shift
of the degradation temperature was reported for
ramie whiskers.
Compared to the neat PEO, the presence of a low
amount of cellulose whiskers does not modify the
thermal stability of composites. For filled films, the
degradation process under helium occurs in two
distinct steps, starting by the cellulose and followed
by that of PEO (this two steps process is clearly
observed for the PEO ? 30 wt% WR). The same
degradation behavior, involving a two steps process,
10
4
10
3
10
2
10
1
10
0
-100
-50
0
50
Temperature (°C)
Fig. 4 Normalized storage tensile modulus E0 for cast/
evaporated unfilled PEO (filled square) and nanocomposite
films reinforced with 3 (open circle), 6 (filled triangle), 10
(inverted triangle), 20 (diamond) and 30 wt% (left-pointing
triangle) of ramie whiskers as a function of temperature
123
Cellulose
The relaxation associated with the glass transition
of the amorphous domains of PEO occurs at about
-55 °C. The modulus drop corresponding to this
relaxation is weak because of the high degree of
crystallinity of PEO. Above -55 °C, E0 decreases
continuously because of the progressive softening of
PEO. When reaching the melting point of the
polymeric matrix around 70 °C, the modulus drops
irreversibly for the neat PEO.
When adding ramie whiskers, the rubbery modulus, observed below the PEO melting temperature,
slightly increases due to the reinforcing effect of
whiskers. Indeed, due to the high crystallinity content
of PEO, the rubbery modulus of PEO is very high and
was weakly modified by whiskers addition. The main
effect is the thermal stabilization of the storage
modulus above the melting point of PEO, Tm, up to
temperatures higher than 100 °C.
As stressed in Table 2, the value of this high
temperature modulus increases as the whiskers content
in the nanocomposite film increases. This phenomenon
may be ascribed to the formation of a rigid percolating
cellulose whiskers network within the polymeric
matrix through strong whiskers/whiskers hydrogen
bonds interaction (Takayanagi et al. 1964). The
modulus of this continuous network can be well
predicted from the adaptation of the percolation
concept to the classical series–parallel model. In this
model and at sufficiently high temperature, i.e. when
the storage modulus of the matrix is much lower than
that of the percolating network, the following equation
was derived (Dufresne et al. 1997) for the predicted
elastic modulus, E0 pre, of the composite:
0
¼ WER0
Epre
ð2Þ
With:
W ¼ 0 for
tR [ tRc
Table 2 High temperature (T = 80 °C) tensile modulus:
comparison between experimental (E0 exp) and predicted (E0 pre)
data for ramie whiskers reinforced PEO nanocomposite films
Sample
3 wt%
6 wt%
10 wt%
20 wt%
30 wt%
E0 exp (MPa)
8
12
20
40
60
E0 pre (MPa)
–
4
9
20
34
123
With:
W ¼ tR
tR tRc
1 tRc
b
for
tR [ tRc
ð3Þ
where W and E0 R are the volume fraction and the elastic
modulus of the rigid percolating network, respectively;
tR, tRc and b correspond to the volume fraction of filler,
critical volume fraction of filler at the percolation
threshold and the corresponding critical exponent,
respectively. For a 3D network, b = 0.4 (De Gennes
1979) and tRc = 2.5 vol% was determined from the
aspect ratio of ramie whiskers, L/d = 28.
The tensile modulus of dry ramie whiskers films,
E0 R, was experimentally determined and a value of
about 0.35 GPa was found. This value results from
the average of two experiments that were relatively
reproducible despite the extreme brittleness of the
films. This brittleness is no more observed in
the PEO/whiskers composite. The low value of the
tensile modulus obtained for ramie whiskers, could
be associated with the low aspect ratio of ramie
whiskers. Indeed, Bras et al. (Bras et al. 2010) shown
a correlation between the tensile modulus and the
aspect ratio for a large number of whisker sources.
For the predicted modulus, the densities of ramie
whiskers and PEO were taken as 1.5 and 1.2 g cm-3,
respectively. The predicted storage modulus values,
E0 pre, are reported in Table 2. They were not determined for the nanocomposite film reinforced with
3 wt% of ramie whiskers, because this filler content is
slightly lower than the theoretical percolation threshold value. However, experimentally, a stabilization of
the storage tensile modulus was observed with 3 wt%
(tRc = 2.41 vol%) whiskers PEO composite, thus
very close to the theoretical percolation threshold.
This difference may be associated with the model
developed, which neglects the effect of whiskers
reinforcement below the percolation threshold.
Regardless the composition of the sample, experimental modulus values, E0 exp, display a similar evolution to the predicted one even if experimental values
were normalized at low temperature. It is a good
indication that the stiffness of the sample and the
temperature stabilization of the composite modulus
most probably result from the formation of an H-bonded
cellulose whiskers network as proposed in the model.
Even if the H-bonded strength decreases with
increasing temperature, the large number of H-bonds
Cellulose
involved in whiskers/whiskers interaction induces the
stabilization of the composite storage modulus (Favier et al. 1997).
A previous study has shown that for composites
based on a PEO matrix and tunicin whiskers, the
experimental high temperature modulus values were
about 18, 45, and 235 MPa for composites filled with
3, 6, and 10 wt% whiskers, respectively (Azizi Samir
et al. 2005b). Compared to ramie whiskers-based
nanocomposites, the higher modulus values obtained
for tunicin whiskers-based nanocomposites are mainly
ascribed to both the higher aspect ratio of tunicin
whiskers, of about 70 (only 30 for ramie whiskers) and
the higher elastic modulus of the tunicin network.
Impact of film processing
Extrusion is an industrial method allowing to manufacture a large range of products in short times. We
investigated the effect of this industrial process on the
properties of ramie whiskers reinforced PEO nanocomposite films. The whiskers content was fixed at
6 wt%, i.e. 4.86 vol%. Indeed, this amount is higher
than the percolation threshold (2.5 vol%) and corresponds to the optimum balance between a low
whiskers content and a strong reinforcing effect.
The processing may have a direct impact on both
the thermal and mechanical properties of the composite films because the extrusion process induces
mechanical and temperature stresses and some possible orientation of the fibers.
In order to determine the optimized extrusion
conditions, the isothermal stability of PEO was
investigated by TGA and rheological measurement
under inert atmosphere. Neat PEO and PEO reinforced with 6 wt% ramie whiskers samples were
maintained at 180 °C for 8 h, and the weight loss
observed was only equal to 2% of the initial sample
weight and was associated to water evaporation. This
result is in accordance with data obtained upon
heating in Table 1 and Fig. 2, which show that the
onset degradation temperature of PEO was well
above 180 °C. The Fig. 5 shows the evolution of both
G0 and G00 moduli versus time at 180 °C and 1 Hz for
the matrix obtained by melting PEO powder. The
moduli are constant indicating the stability of PEO at
180 °C in inert atmosphere. The extrusion process
was thus performed in a twin screw, under nitrogen
flow in order to avoid the oxidation of PEO at
Fig. 5 Storage (G0 , filled square) and loss (G00 , open circle)
moduli versus time at 180 °C for neat PEO, frequency 1 Hz,
deformation amplitude 0.5%
180 °C. This temperature seems to be a good
compromise between low viscosity and thermal
stability. The extrusion speed was maintained at a
low value, i.e. 25 rpm, to limit the PEO chains and
whiskers break.
Morphology of extruded films
The morphology of the extruded nanocomposite film
reinforced with 6 wt% of ramie whiskers was characterized by SEM. Figure 6 shows the cryofractured
surface of this material.
The morphology of the extruded nanocomposite
film is similar but less chaotic than its cast/evaporated
counterpart (Fig. 1b). The extruded film didn’t
Fig. 6 Scanning electron micrographs of cryofractured surface
of the extruded nanocomposite film reinforced with 6 wt%
ramie whiskers
123
Cellulose
significant narrowing of the length distribution,
showing that ramie whiskers are more monodisperse
in length after extrusion. Such modifications in the
length distribution can be properly attributed to a
more efficient degradation of longer whiskers.
For this extruded composite, the cross section and
length of whiskers were averaged over 300 measurements and they were found to be 5 ± 1 and
122 ± 45 nm, respectively, giving an aspect ratio
around 24 ± 17. These values were compared to the
initial average values of 7 ± 1, 200 ± 78 and
28 ± 12 nm, respectively. These results show that
the extrusion process do not induce a significant
change of the aspect ratio of the rod-like cellulosic
nanoparticles. Indeed, even if individual variations of
the cross section and length were reported, the impact
on both is mostly equivalent.
display voids but large domains of white dots
indicating that the whiskers are not well dispersed
in the PEO matrix. The dots observed are much larger
than those obtained for the cast/evaporated film,
shown in Fig. 1b. The freeze-dried sample, before
extrusion, does not exhibit such morphology. Consequently the whisker aggregates may have been
induced by the extrusion process.
In order to evaluate the influence of the extrusion
process on the whiskers degradation, the whiskers
length and diameter after extrusion were determined
through TEM observations. Ramie whiskers were
extracted from the extruded composite material by
dissolving the composite in water. After dissolution
of the PEO matrix, the cellulose whiskers were
observed by TEM and compared to that directly
obtained from the aqueous suspension (Fig. 7a, b).
Uranyl acetate at a concentration of 2 wt% was used
in order to emphasize the cellulose whiskers and
create contrast.
As stressed in Fig. 7c, the effect of the extrusion
process on the length of ramie whiskers is twofold:
the extrusion greatly decreases the length of the main
population, characterized by the peak position in the
distribution, by a factor of about two, passing from
about 200 to 120 nm. The second effect is the
Fig. 7 Transmission
electron micrographs
(TEM) of a ramie whiskers
suspension; b extruded and
re-dispersed PEO
nanocomposite films
reinforced with 6 wt% of
ramie whiskers; c their
length distributions
Rheometry
The rheometrical characterization of PEO-based composites was performed through viscoelastic and creep
measurements. For viscoelastic measurements, the
linear regime was previously determined for each
sample through a strain sweep test. At 90 °C, i.e. above
the melting temperature, the critical strain, cc, marking
(a)
(b)
200nm
Numbers of Whiskers
(c)
Whiskers before Extrusion
Whiskers after Extrusion
60
50
40
30
20
10
0
0
50
100
150
200
Length (nm)
123
250
300
Cellulose
(a)
5
10
G'; G" (Pa)
the upper limit of the linear regime was about 0.5% for
the cast/evaporated and extruded unfilled matrix while
it decreased to 0.08% for the nanocomposites. The
decrease of cc generally observed for encumbered
systems, is due to the presence of ramie whiskers for
nanocomposites.
Figure 8a, b show the evolution of both G0 and G00
moduli as a function of the angular frequency for the
neat PEO films and nanocomposites filled with
6 wt% of ramie whiskers, obtained by casting/evaporation and extrusion, respectively. The complex
viscosity for all materials is presented in Fig. 8c.
4
10
0.4
ω
3
10
0.8
ω
G'
G"
2
10
(b)
5
10
4
10
3
10
G'
G"
2
10
6
10
5
10
η * (Pa.s)
Let’s first examine the viscoelastic response of the
matrix, processed by either casting/evaporation or
extrusion.
In the case of cast/evaporated PEO films, the
viscoelastic behavior is typical of melt polymers with
the onset of a terminal zone at low frequency and the
beginning of a rubbery plateau at high frequency,
separated by a G0 –G00 cross over at intermediate
frequency. It has to be stressed that the frequency
dependences of both moduli in the terminal zone, i.e.
G0 µ x0.8 and G00 µ x0.4, are quite lower than those
expected for dense molecular systems with exponents
of 2 and 1, respectively. Such lower frequency
dependence of viscoelastic moduli in the terminal
zone has been observed for PEO solutions and has
been attributed to the presence of aggregates (Bossard
et al. 2010). In the bulk, it could be the rheological
signature of the presence of crystallites or spherulites
in the amorphous phase.
The extruded PEO film (Fig. 8b) exhibits a viscoelastic behavior similar to the one of the cast/
evaporated film with some quantitative differences:
(1) the levels of both moduli, and consequently the
complex viscosity, are lower for the extruded polymer
than for the cast/evaporated one and (2) the terminal
zone is shifted towards very low frequencies not
explored and the G0 –G00 cross-over is shifted towards
higher frequencies. For the latter effect, it points out
that the average relaxation time dynamics k of PEO
molecules, corresponding roughly to the inverse of the
frequency of G0 –G00 crossover, is speeded up after
extrusion, passing from 2 to 0.2 s.
A decrease in the viscosity, associated with a speed
up of the molecular dynamics and the broadening of
G'; G" (Pa)
Matrix behavior
4
10
3
(c)
10
-3
10
-2
10
-1
10
0
10
1
10
2
10
ω (rad/s)
Fig. 8 Storage (G0 , filled symbols) and loss (G00 , open symbols)
moduli versus angular frequency at 90 °C for a the cast/
evaporated unfilled PEO matrix (circle) and 6 wt% ramie
whiskers reinforced nanocomposites (square), and b the
extruded unfilled PEO matrix (triangle) and 6 wt% ramie
whiskers reinforced nanocomposites (square). c Complex
viscosity of the cast/evaporated (square) and extruded films
(triangle) for matrices (open symbols) and composites (filled
symbols) versus angular frequency
the frequency region between the terminal zone and
the G0 –G00 cross-over after extrusion could be ascribed
to a chain scission effect with a broadening in the
polydispersity index of the polymer. Indeed, a similar
effect has been observed for stirred PEO water
solutions, and attributed mainly to the elongational
flow induced by the dispersion procedure (Bossard
123
Cellulose
et al. 2010). Under extrusion that induces intense
elongational flow, polymer chain scission is highly
expected. To confirm this hypothesis, the average
molecular weight of the extruded polymer was compared to the one obtained for the polymer processed by
casting/evaporation using viscosity measurements.
For this purpose, both films were dissolved and diluted
at several concentrations in distilled water. The
intrinsic viscosity [g] was determined as the extrapolation to zero concentration of the reduced viscosity
gred. defined in Eq. 4
gred: ¼
gs gw
Cgw
ð4Þ
with C, the solution concentration, gs the zero-shear
viscosity of polymer solutions and gw = 0.97 mPas
the Newtonian viscosity of water at 21 °C. Alternatively, [g] can be obtained by fitting the so-called
inherent viscosity, ginh = (ln grel)/c with the Kraemer
equation
ln grel
¼ ½g kK ½g2 c
c
ð5Þ
where grel is the relative viscosity, grel = g0/gw and
kK the Kraemer coefficient. The intrinsic viscosity is
directly related to the molecular weight M by the
Houwink-Mark-Sakurada equation (HMS),
½g ¼ KM a
ð6Þ
where K and a are constants (Flory 1953). For PEO,
HMS constants at 25 °C are equal to K =
6.103 9 10-3 cm3 g-1 and a = 0.83 (Khan 2006).
The intrinsic viscosity of PEO solutions obtained
from the cast/evaporated film is about [g]cast–evap =
900 ± 150 cm-3 g-1 while the one measured for the
extruded film is [g]cast–evap = 520 ± 80 cm-3 g-1,
corresponding to an average molecular weight
Mcast–evap = (1.69 ± 0.2) 9 106 g/mol and Mextr. =
(8.7 ± 1.6) 9 105 g/mol, respectively. It thus
appears that the decrease of the viscosity is effectively due to the significant mechanical degradation
of PEO molecules after extrusion through chain
scission.
Composite behavior
Let us consider and compare now the viscoelastic
behavior of the two nanocomposites obtained either
by casting/evaporation or extrusion. It can be seen in
123
Fig. 8a, b that both materials exhibit viscoelastic
moduli and a complex viscosity higher than that of
their respective matrices, confirming the mechanical
strengthen induced by the whiskers via whiskers/
whiskers and whiskers/PEO interactions.
A similar behavior was reported by (Alvarez et al.
2004), with a saturation effect at higher fibers
content. Indeed, strong interactions exist between
PEO chains and cellulose. This effect is exacerbated
because of the large cellulosic surface inherent to any
nanoparticle. The PEO molecular dynamic is therefore locally restricted in the interfacial regions. This
result is consistent with the pulse field NMR studies
reported by (Azizi Samir et al. 2004b) for tunicin
whiskers/PEO nanocomposites. These authors
showed that the long relaxation time of PEO chains
strongly decreased even with a low amount of
whiskers owing to whiskers/PEO chains interactions.
However, some significant differences can be
noticed between cast/evaporated and extruded nanocomposites. Viscoelastic moduli for the cast/evaporated nanocomposite in Fig. 8a are nearly frequency
independent with G0 [ G00 , except at very low
frequency. This solid-like behavior would suggest
the presence of a physical network, probably composed of ramie whiskers.
For the extruded nanocomposite, the viscoelastic
moduli are frequency dependent and slightly lower
than those of the cast/evaporated nanocomposite.
After extrusion, the viscoelastic behavior is frequency depended thus suggesting the absence of a
network, contrarily to what was observed for the cast/
evaporated nanocomposite.
Consequently, it can be supposed that the extrusion prevents the formation of a network. In order to
verify this hypothesis, creep measurements obtained
for cast/evaporated and extruded nanocomposites
submitted to the same stress have been compared in
Fig. 9. The strain of the cast/evaporated nanocomposite reaches a plateau value beyond 1,500 s,
corresponding to the mechanical response of a solid
with a delayed elasticity while the one of the extruded
nanocomposite gradually increases with time, which
is characteristic of a fluid.
These results confirm the formation of a network
for the cast/evaporated nanocomposite. In the case of
the extruded nanocomposite, the whole set of rheological data suggests that whiskers do not form a
network in the extruded film. However, the formation
Cellulose
–
7
Extruded
6
–
Strain (%)
5
4
–
3
2
Cast/evaporated
1
–
0
0
1000
2000
3000
4000
Time (s)
Fig. 9 Creep measurements (s = 5 lNm) for extruded (open
circle) and cast/evaporated (filled square) composite reinforced
with 6 wt% of ramie whiskers at 90 °C under inert atmosphere
of a weak network through low density H-bonds
cannot be excluded.
Let us compare now the average relaxation time k
corresponding to the inverse of the frequency at the
G0 –G00 cross over.
In the case of PEO films, k is divided by a factor of
about 10 after extrusion. This speed up in the
molecular dynamic has been attributed to PEO chain
scission. For nanocomposites, the average relaxation
time is divided by a factor of about 36, passing from
180 to 5 s.
As a consequence, differences in the relaxation
dynamics of nanocomposites cannot be explained
only by the modification of the matrix after extrusion
but could be also due to the combined mechanical
degradation and aggregation of whiskers, as stressed
in Fig. 6 by SEM investigation and in Fig. 7a, b by
TEM measurements. Indeed, break up of cellulose
whiskers (Pathi and Jayaraman 2006) or natural fibers
(Bengtsson et al. 2007) upon extrusion were already
reported. (Alvarez et al. 2004) have shown that
rheological properties of composite material are very
sensitive to the diameter and aspect ratio of the fibers.
Any process inducing a decrease of the fiber cross
section and aspect ratio results in lower viscosity
values.
Consequently, from a microstructural point of
view, the general decrease of the rheological properties of the extruded nanocomposites compared to
cast/evaporated ones may likely result from the
contribution of four combined effects:
The decrease of the rheological properties of the
matrix through PEO chain scission induced by
extrusion.
The mechanical degradation of ramie whiskers
during extrusion that reduces the ability of
cellulosic fibers to connect each other.
The whiskers aggregation induced by the extrusion process, which decreases the amount of
whiskers available for the formation of the
percolating network, as reported from SEM
observation.
And also the expected orientation effect of the
extrusion process that prevents the formation of
the percolation network.
Thermal characterization
Non-isothermal investigation
The thermal behavior of extruded samples was
characterized using DSC. Results are reported in
Table 3. The extruded PEO matrices obtained using
either the PEO powder or pellets of freeze-dried PEO
solution present similar thermal properties. Compared to the cast/evaporated neat sample (Table 3),
the extruded neat samples display similar crystallization temperatures of about 51 °C, while both their
melting temperatures and degrees of crystallinity
decrease. A low value of the melting point is
generally associated with a low value of lamellar
thickness or/and a high value of the end interfacial
free energy. These two parameters strongly depend
on the crystallization process, i.e. melt crystallization
or polymer precipitation in solution.
Table 3 Thermal characteristics of extruded PEO-based
nanocomposites reinforced with ramie whiskers obtained from
DSC curves: crystallization temperature (Tc), glass transition
temperature (Tg), and degree of crystallinity expressed as a
function of the matrix weight measured during the melting (vm)
Samples
Tg (°C) Tc (°C) Tm (°C) vam
Extruded PEO powder
-55
52
49
0.76
Extruded freeze-dried PEO
-55
49
50
0.75
Extruded PEO ? 6 wt% WR -53
51
41
0.7
mPEO
vm ¼ DHDH
where
= 210 J g
0
m
for 100% crystalline PEO
a
DH0m
-1
is the heat of melting
123
Cellulose
50°C
800
Average radius (nm)
The significant differences in Tm and vm may be
ascribed to (1) the processing technique itself, the
spherulites size depending on the crystallization
conditions, i.e. from the polymer melt or by polymer
precipitation in solution, (2) the polymer chain
scission during extrusion which may involve some
polymer ramification and (3) the polymer orientation.
Because of the lower degree of crystallinity, the glass
transition can be observed at about -55 °C for
extruded PEO matrices.
For the extruded samples, the Tm value significantly decreases, with a difference of 8 °C between
filled and unfilled samples. This effect is much higher
than the one observed for the cast/evaporated films
for which no decrease of the melting point was
observed with the incorporation of 6 wt% whiskers.
The crystallization temperature was found to remain
roughly constant.
600
400
55°C
200
0
50
100
150
200
250
300
350
Time (s)
Fig. 10 Time dependence of the spherulites radius of extruded
PEO matrix (filled circle) and cast/evaporated PEO matrix
(open square) at 50 °C and extruded PEO matrix (inverted
triangle) and extruded composite with 6 wt% of ramie
whiskers (open triangle) and cast/evaporated PEO composite
with 6 wt% of ramie whiskers (diamond) at 55 °C
Isothermal crystallization
In order to try to elucidate the crystallization process,
isothermal crystallization kinetics were investigated.
The growth rate of the PEO spherulites was determined for cast/evaporated and extruded matrix films
using polarized optical microscopy. The sample was
first melt at 100 °C for few minutes and cooled down
to 50 or 55 °C. The linear growth rate is very
sensitive to the imposed crystallization temperature.
Indeed, due to the large difference in the behavior of
the materials studied, two crystallization temperatures were investigated in order to crystallize the
material in appropriate time.
For both matrices, cast/evaporated and extruded
ones, the evolution of the PEO spherulites radius was
monitored at 50 °C and results are reported in Fig. 10.
The spherulites radii increase linearly with time for the
two samples, which is generally observed for isothermal polymer crystallization. The kinetic of the radius
growth is similar. However, a large difference exists,
associated with the number of spherulites formed at a
time t. The germ density, as observed during the optical
investigation, for the extruded sample is high, thus the
coalescence of spherulites occurs quickly and avoids
the measurement of their radii beyond 50 s. As the
germ density of the cast/evaporated matrix is much
lower than for extruded one, the total crystallization is
obtained after 300 s with a low density of large
spherulites. The final spherulites radius sizes were
123
estimated around 850 and 260 nm for the cast/evaporated and extruded matrices, respectively.
The linear growth rate of spherulites for extruded
PEO and nanocomposites has been studied at 55 °C.
The increase of the crystallization temperature
involves the presence of an induction time, necessary
to obtain the first spherulites germ. The extruded
composite sample exhibits the same linear growth
rate of spherulites than the extruded PEO matrix.
However, the setup fails to access the spherulites
cross section for extruded PEO because it stops
rapidly, after 180 s, due to the coalescence of the
spherulites, as observed at 50 °C. The two composite
samples exhibit the same kinetic, with the coalescence of the spherulites obtained after 300 s. For both
extruded and cast/evaporated composites, the final
spherulites radius was estimated around 330 nm.
Therefore, it appears that the elaboration process
has an effect on the isothermal crystallization for the
neat samples by a modification of the germ density.
The extrusion process induces PEO degradation with
a significant molecular weight decrease. This degradation may induce some defects, i.e. chain ramification which may increase the germ density as observed
during optical measurements.
The incorporation of whiskers seems to vanish the
influence of the processing technique on the linear
growth rate of spherulites and their final radius. This
may be related to the large influence of the presence
Cellulose
Thermal degradation
The thermal degradation of the unfilled PEO and
composite films was investigated under air. Composites and unfilled extruded films present a lower thermal
stability than the cast/evaporated ones. Composites
and unfilled extruded films have an onset degradation
temperature of about 191 and 178 °C, respectively,
compared to 205 and 191 °C for unfilled and composite cast/evaporated samples. The degradation process
of PEO in the presence of oxygen is very complex
(Costa et al. 1992). The lower degradation temperature
of extruded films may be explained by the fact that
during extrusion a significant mechanical degradation
of PEO molecules through chain scission occurs. It
may involve the formation of weaker links or end
groups which are more sensitive to oxidative thermal
degradation. Under helium, no effect of the processing
technique was observed on thermal degradation and
this invariance may be associated to the less aggressive
atmosphere for PEO degradation.
Mechanical behavior
Dynamic mechanical measurements were performed
for the extruded samples and compared to those
obtained for the cast/evaporated films in Fig. 11.
Here again, the storage tensile modulus E0 at
-100 °C was normalized at 8 GPa to minimize the
influence of the error made for the determination of
the sample dimensions. In accordance with its lower
degree of crystallinity, as revealed by DSC measurements, the extruded PEO exhibits a higher modulus
drop at Tg compared to the cast/evaporated matrix
(Fig. 11a). Then, after the glass transition, the
modulus continuously decreases because of the
progressive melting of crystalline domains of PEO
up to the melting point. The storage tensile modulus
decrease near 65 °C is similar for the two matrices.
Dynamic mechanical measurements were performed for extruded nanocomposite films reinforced
(a)
4
10
3
10
Normalized storage modulus (MPa)
of the whiskers on the crystallization process. The
incorporation of 6 wt% of whiskers has no effect on
the linear growth rate, and thus doesn’t modify the
polymer chain mobility involved in the crystallization
process. For cast/evaporated samples, the incorporation of 6 wt% whiskers involves an increase of the
nucleation density.
2
10
1
10
4
10
(b)
3
10
2
10
1
10
-100
-50
0
50
100
150
Temperature (°C)
Fig. 11 Normalized storage modulus E0 for a the neat cast/
evaporated (open circle) and extruded (filled square) PEO
matrices, and b nanocomposite films reinforced with 6 wt% of
ramie whiskers obtained by casting-evaporation (filled triangle), strained in the extruded direction (open circle) and
strained in the cross-sectional direction (filled square) as a
function of temperature
with 6 wt% of ramie whiskers (Fig. 11b) for samples
cut in the extrusion direction and in the crosssectional one. (Kvien and Oksman 2007) have shown
that using a strong magnetic field, which induces a
cellulose whiskers orientation, the nanocomposite
film modulus in the cross-sectional direction is higher
than that in the extrusion one. For extruded PEO
nanocomposites, the curves obtained for the sample
cut in the two directions are overlapped indicating
that no orientation of the whiskers occurs during
extrusion. Thus one hypothesis developed in regard
to rheological measurements to explain the decrease
in rheological properties is suppressed by mechanical
investigation.
The experimental modulus of the extruded composite was found around E0 exp = 2 MPa, compared to
12 MPa for the cast/evaporated nanocomposite film
reinforced with 6 wt% of ramie whiskers. The high
temperature storage modulus of the filled extruded
123
Cellulose
membrane is therefore notably lower than the one of
the filled cast/evaporated membrane.
Even if the whisker aspect ratio was slightly
reduced as discussed previously, 24 instead of 28, the
amount of whiskers in the extruded membrane, i.e.
6 wt% (4.86 vol%), is theoretically sufficient to
obtain a percolating network because the critical
volume fraction of the filler at the percolation
threshold calculated using an aspect ratio of 24 is
tRc = 2.9 vol%. Nevertheless, the presence of whisker aggregates in the filled extruded membrane,
observed by MEB investigation, decreases notably
the amount of whiskers available to form a network.
Moreover, in cast/evaporated films, the kinetic of the
film formation and the viscosity of the medium, at
least at the beginning of the process are low. These
two parameters give time to the formation of a
whiskers network in cast/evaporated membranes.
On the contrary, during the extrusion process, a
high viscosity value for the matrix, a fast process
kinetic and mechanical stress could restrict the
number of H-bonds formed during whiskers network
formation. Thus a weak network or in the limit case
no network may be formed.
At high temperature, instead of the E0 exp plateau,
obtained in the case of cast/evaporated films characterized by DMA measurement (Fig. 11), a straight
line with a slope of about -0.03 MPa °C-1 for
extruded composites is observed. This behavior may
be associated with the absence of network or the
presence of a weak network.
As mentioned previously, the invariance of E0 exp
with temperature for cast/evaporated composites was
attributed to a high density of H-bonds. Thus the
decrease of E0 exp versus temperature for extruded
sample is in agreement with a low density H-bonded
network. Indeed, as the H-bonded strength decreases
with temperature, a low density H-bonded network
may have the same behavior.
Conclusions
Nanocomposite films based on PEO polymer as the
matrix and cellulose whiskers extracted from ramie
plant as the reinforcing phase were obtained by
casting/evaporation and extrusion processes. Microscopic observations show some whiskers aggregations and a small decrease of the whiskers aspect
123
ratio for extruded sample, but for both processes
employed, films display homogeneous surfaces.
The rheological behavior for cast/evaporated films
shows that viscoelastic and creep measurements have
a solid-like behavior, according to mechanical measurement exhibiting a spectacular reinforcement after
melting temperature. These high mechanical performances for the casting/evaporation process are
ascribed to the formation of a rigid cellulosic
network.
For the extruded composites, the rheological
behavior through the viscoelastic and creep measurements shows a liquid-like behavior. This stresses a
weak reinforcement behavior obtained for extruded
composites. This weak mechanical reinforcement
after the PEO melting temperature leads us to
conclude that the extrusion process prevents the
formation of a strong whiskers network for the
whiskers content used, contrary to evaporated films.
One possibility to improve this reinforcement, taking
into account the aggregation and the decrease of the
aspect ratio of whiskers during the extrusion process,
is to increase the filler content. Even if the extrusion
process requires a higher whisker content to obtain
the same result compared to the casting/evaporation
process, it is worth noting that the former is much less
time consuming.
Acknowledgments The authors thank Dr. Youssef Habibi for
his support in the whiskers preparation and Mme Denise
Foscallo for TGA measurements.
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