Polímeros: Ciência e Tecnologia (Polimeros) 4th. issue, vol 27, 2017

DRIVING INNOVATION FORWARD Polímeros

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VOLUME XXVII - Issue IV - OCT/DEC - 2017

October 22-26, 2017 Hotel Majestic Convention Center Águas de Lindóia, SP, Brazil © 2016 Copyright SABIC. All rights reserved.

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Redução no consumo de solvente e consequentemente na geração de resíduos Novas tecnologias de colunas que permitem rápida troca de solvente

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http://dx.doi.org/10.1590/0104-1428.2704

Polímeros has a new President of the Council Board and two newly appointed Associate Editors 2017 has been a special and productive year to Polímeros. Last October, 24 during the 14th Congresso Brasileiro de Polímeros, which was held in Águas de Lindóia, SP, the members of the Council Board met and various important resolutions were taken. One of them was the appointment of the new President of the Council Board, Prof. Antonio Aprigio da Silva Curvelo, from Instituto de Química de São Carlos, Universidade de São Paulo, São Carlos, SP. He substitutes Prof.

Marco-Aurelio de Paoli, from Instituto de Química, Universidade de Campinas, SP who is taking the Vice-Presidency of Associação Brasileira de Polímeros, ABPol. We acknowledge the great work Prof. de Paoli has done to this journal during his long staying and salute Prof. Curvelo to his new position. Another important decision was the appointment of two new members, Prof. Paula Moldenaers from University of Leuven, Belgium and Prof. Sadhan C. Jana from University of Akron, USA (below a short CV of both). They were at the same meeting, appointed as new international Associated Editors, joining Profs. Alain Dufresne (Grenoble INP/Pagora, France), José A. Covas (UMinho/IPC, Portugal) and Richard G. Weiss (GU/DeptChemistry, USA). We all welcome the new President and the members of the Associate Editors.

From January, 2017 the Council Board have approved the article-charge and it was promptly implemented to all accepted articles submitted to Polímeros since then. At the beginning a few authors of accepted articles disagreed with this helping fund. Today, just after 3 issues being published, all authors of accepted article did agree to contribute funding the journal. All members of the Council and Editorial team acknowledge this assertiveness, sign of a well-established community.

This year Polímeros has published 64 articles in all, 48 in its four standard issues. An extra issue containing the very last 16 accepted articles written in Portuguese was published in the beginning of this year. So far the journal had a submission of 86 articles, from which 40 were accepted. It is noteworthy to mention the increase of the submissions coming from the Brazilian states of Paraná and Paraíba. Submissions from abroad are also in continuous increase, especially from Colombia, India, Iran and Turkey. Terms like emulsion polymerization and biodegradable polymers were the most common search keywords used by the visitors of the journal’s webpage.

For all that I alleged that 2017 has being a particularly productive year for Polímeros.

Sebastião V. Canevarolo Editor-in-Chief Short CV of the newly appointed President and Associate Editors of Polímeros.

President of the Council Board Prof. Antonio Aprigio S. Curvelo from Institute of Chemistry of São Carlos, University of São Paulo, São Carlos, Brazil. His scientific work is dedicated to vegetal macromolecules, in particular dedicated to the separation of macromolecules from lignocellulosic biomass by using organic solvents and fluids in the sub/supercritical states.

His work also includes the characterization of cellulose, hemicelluloses and lignins, and the production of micro‑and macromolecular derivatives and composite materials in a context of Biorefinery.

Associate Editor Prof. Paula Moldenaers, from Department of Chemical Engineering, University of Leuven, Leuven, Belgium.

Her research interests focus on the rheology and morphology development in complex polymeric systems, especially two‑phase polymeric blends, filled polymers and biopolymeric systems. The research aims at providing a science‑based methodology for the processing of such complex fluids, developing and using advanced experimental approaches which have resulted in unique experimental set-ups, data sets, characterization procedures, scaling relations and morphological insight.

Associate Editor Prof. Sadhan C. Jana from Department of Polymer Engineering, University of Akron, Akron, USA. His research interests are mainly centered in understanding shape memory polymers and their nanocomposites and applications.

Also he is interested in the design of nanocomposites based on polyolefins, and the development of mechanisms of nanoparticle-induced morphology in immiscible polymer systems.

Polímeros, 27(4), 2017

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E E E E E E E E E E E E E E E E E E E E E E E E E

ISSN 0104-1428 (printed)

E E E E E E E E E E E E E E E E E E E E E E E E E E E E

ISSN 1678-5169 (online)

P o l í m e r o s - I ss u e I V - V o l u m e X X V I I - 2 0 1 7 I n d e x e d i n : “ C h e m ic a l A b s t r a c t s ” — “ RA P RA A b s t r a c t s ” — “A l l - R u s s i a n I n s t i t u t e o f S ci e n c e ­T e c h n ic a l I n f o r m a t i o n ” — “ R e d d e R e v i s t a s C i e n t i f ic a s d e A m e r ic a L a t i n a y e l C a r i b e ” — “ L a t i n d e x ” — “ W e b o f S ci e n c e ”

and

Polímeros E d i t o r i a l C o u nci l

Editorial Committee

Antonio Aprigio S. Curvelo (USP/IQSC) - President

Sebastião V. Canevarolo Jr. – Editor-in-Chief

Members Adhemar C. Ruvolo Filho (UFSCar/DQ) Ailton S. Gomes (UFRJ/IMA) Alain Dufresne (Grenoble INP/Pagora) Antonio Aprigio S. Curvelo (USP/IQSC) Bluma G. Soares (UFRJ/IMA) César Liberato Petzhold (UFRGS/IQ) Cristina T. Andrade (UFRJ/IMA) Edson R. Simielli (Simielli - Soluções em Polímeros) Edvani Curti Muniz (UEM/DQI) Elias Hage Jr. (UFSCar/DEMa) Eloisa B. Mano (UFRJ/IMA) João B. P. Soares (UAlberta/DCME) José Alexandrino de Sousa (UFSCar/DEMa) José António C. Gomes Covas (UMinho/IPC) José Carlos C. S. Pinto (UFRJ/COPPE) Júlio Harada (Harada Hajime Machado Consutoria Ltda) Laura H. de Carvalho (UFCG/DEMa) Luiz Antonio Pessan (UFSCar/DEMa) Luiz Henrique C. Mattoso (EMBRAPA) Marco-Aurelio De Paoli (UNICAMP/IQ) Osvaldo N. Oliveira Jr. (USP/IFSC) Paula Moldenaers (KU Leuven/CIT) Raquel S. Mauler (UFRGS/IQ) Regina Célia R. Nunes (UFRJ/IMA) Richard G. Weiss (GU/DeptChemistry) Rodrigo Lambert Oréfice (UFMG/DEMET) Sadhan C. Jana (UAKRON/DPE) Sebastião V. Canevarolo Jr. (UFSCar/DEMa) Silvio Manrich (UFSCar/DEMa)

A ss o ci at e E d i t o r s Adhemar C. Ruvolo Filho Alain Dufresne Bluma G. Soares César Liberato Petzhold José António C. Gomes Covas José Carlos C. S. Pinto Paula Moldenaers Regina Célia R. Nunes Richard G. Weiss Rodrigo Lambert Oréfice

Sadhan C. Jana

D e s k t o p P u b l is h in g

www.editoracubo.com.br

“Polímeros” is a publication of the Associação Brasileira de Polímeros São Paulo 994 St. São Carlos, SP, Brazil, 13560-340 Phone: +55 16 3374-3949 emails: abpol@abpol.org.br / revista@abpol.org.br http://www.abpol.org.br Date of publication: December 2017

Financial support:

Polímeros / Associação Brasileira de Polímeros. vol. 1, nº 1 (1991) -.- São Carlos: ABPol, 1991Available online at: www.scielo.br

Quarterly v. 27, nº 4 (Oct./Nov./Dec. 2017) ISSN 0104-1428 ISSN 1678-5169 (electronic version)

Website of the “Polímeros”: www.revistapolimeros.org.br

1. Polímeros. l. Associação Brasileira de Polímeros.

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Editorial Section Editorial................................................................................................................................................................................................E1 News....................................................................................................................................................................................................E4 Agenda.................................................................................................................................................................................................E5 Funding Institutions.............................................................................................................................................................................E6

O r i g in a l A r t ic l e Nanostructured magnetic alginate composites for biomedical applications Pedro Marins Bedê, Marcelo Henrique Prado da Silva, André Bem-Hur da Silva Figueiredo and Priscilla Vanessa Finotelli.................... 267

Green polyurethane synthesis by emulsion technique: a magnetic composite for oil spill removal Raphael Maria Dias da Costa, Gabriela Hungerbühler, Thiago Saraiva, Gabriel De Jong, Rafael Silva Moraes, Evandro Gonçalves Furtado, Fabrício Machado Silva, Geiza Esperandio de Oliveira, Luciana Spinelli Ferreira and Fernando Gomes de Souza Junior................................................................................................................................................................... 273

Thermal degradation of polymer systems having liquid crystalline oligoester segment Renato Matroniani and Shu Hui Wang............................................................................................................................................................ 280

New technologies from the bioworld: selection of biopolymer-producing microalgae Roberta Guimarães Martins, Igor Severo Gonçalves, Michele Greque de Morais and Jorge Alberto Vieira Costa...................................... 285

Microalgae biopeptides applied in nanofibers for the development of active packaging Carolina Ferrer Gonçalves, Daiane Angelica Schmatz, Lívia da Silva Uebel, Suelen Goettems Kuntzler, Jorge Alberto Vieira Costa, Karine Rigon Zimmer and Michele Greque de Morais................................................................................................................................... 290

Modification of thermal and rheological behavior of asphalt binder by the addition of an ethylene-methyl acrylate-glycidyl methacrylate terpolymer and polyphosphoric acid Gerson da Silva Pereira and Ana Rita Morales............................................................................................................................................... 298

Dynamic mechanical analysis and crystalline analysis of hemp fiber reinforced cellulose filled epoxy composite Anand Palanivel, Anbumalar Veerabathiran, Rajesh Duruvasalu, Saranraj Iyyanar and Ramesh Velumayil................................................ 309

Production of biodegradable starch nanocomposites using cellulose nanocrystals extracted from coconut fibers Jamile Costa Cerqueira, Josenai da Silva Penha, Roseane Santos Oliveira, Lilian Lefol Nani Guarieiro, Pollyana da Silva Melo, Josiane Dantas Viana and Bruna Aparecida Souza Machado......................................................................................................................... 320

Layer-by-Layer technique employed to construct multitask interfaces in polymer composites Luísa Sá Vitorino and Rodrigo Lambert Oréfice............................................................................................................................................. 330

Influence of tribological test on the global conversion of natural composites Carlos Eduardo Correa, Robin Zuluaga, Cristina Castro, Santiago Betancourt, Analía Vázquez and Piedad Gañán.................................. 339

Effect of concentrations of plasticizers on the sol-gel properties developed from alkoxides precursors Sandra Raquel Kunst, Marielen Longhi, Lilian Vanessa Rossa Beltrami, Lucas Pandolphi Zini, Rosiana Boniatti, Henrique Ribeiro Piaggio Cardoso, Maria Rita Ortega Vega and Célia de Fraga Malfatti........................................................................... 346

Evaluation of Out-of-Autoclave (OOA) epoxy system Fernanda Guilherme, Silvana Navarro Cassu, Milton Faria Diniz, Tanila Penteado de Faria Gonzales Leal, Natália Beck Sanches and Rita de Cássia Lazzarini Dutra........................................................................................................................................................................ 353

Cover: 14th CBPol Arts by Editora Cubo.

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Flexible Capacitor from Dielectric Polymers Improve Capacitors with better performance and improved properties are highly desirable for the development of the next generation of flexible miniaturized personal electronics and pulsed power systems. However, the rapid progress in this field is impeded by limited compatibility of existing materials and preferred processing methods. In a recent paper published in the Journal of Polymer Science: Polymer Physics, Qun-Dong Shen and his colleagues report a promising solution to this problem by blending a small amount of aromatic polythiourea (PTU) into poly(vinylidene fluoride-chlorotrifluoroethylene) P(VDF-CTFE) to boost the reserve capacity and processability of the composite material.

Urea-based linear aromatic dielectric polymers are amorphous polymers with large dipole moments.

These polymers have a very high charge–discharge efficiency or low energy loss. Unfortunately, these dielectric polymers are barely soluble in common polar solvents and they are brittle. In addition, they are incompatible with the widely employed thin-film processing techniques, including solution-casting, hot-pressing, and biaxial orientation.

VDF-based copolymers with chlorotrifluoroethylene (CTFE) and hexafluoropropylene (HFP) have high dielectric constants and large energy densities at high electric field. However, a strong dipole interaction in VDF-based polymers leads to notable polarization hysteresis. As a consequence, the charge–discharge efficiency is low while dielectric loss is large. In order to increase the energy density and the charge–discharge efficiency of the dielectric polymers, inorganic materials are embedded into the polymer matrix, including nanoparticles, ultra-thin nanosheets, and nanofibers.

Dr. Shen and co-authors added a small amount of aromatic PTU into the P(VDF-CTFE) matrix. The aromatic urea-based polymer effectively prevents the early polarization saturation at a field lower than the breakdown field. Meanwhile, PTU brings structural defects into the P(VDF-CTFE) crystals, and reduces composite crystallinity. This research reveals that nonpolar phases, including amorphous regions and paraelectric crystals, play an important role in the energy storage and release behavior of PVDF and its copolymers.

Source: ADVANCED SCIENCE NEWS - http://www. advancedsciencenews.com

Growing Industrialization will influence the Bio plastics and Biopolymers Market Growth 2017‑2025 Bio plastics and Biopolymers are defined as the type of plastics and polymers which are obtained from renewable biomass source such as vegetable oils and fats, starch corn, technology that has an intrinsic network, intelligent control system and home automation system as its basic components. This technology is equipped with home appliances, digital devices and other components that are interconnected and are accessed with the help of a smart gadget from a far off location. The Bio plastics and Biopolymers market has been segmented by product type, by end user industry and geography.

In terms of product type the Bio plastics and Biopolymers market has been segmented into Bio-PET, Bio-PE, PLA, PHA, Starch Blends, Biodegradable polyster and Regenerated Cellulose among others.

Moreover, in terms of end user industry, the Bio plastics and Biopolymers market has been segmented into packaging industry, bottling industry, agricultural industry, automotive industry and consumer products among others. Furthermore, in terms of geography, the Bio plastics and Biopolymers market has been segmented into five geographies namely North America, Europe, Asia Pacific (APAC), Middle East and Africa (MEA) and Latin America.

Increased usage of Bio plastics and Biopolymers as consumer products and increase in investment on agricultural industry by the key vendors across the globe has contributed to the growth of Bio plastics and Biopolymers market. Moreover, the concept of smart cities and smart buildings and connected cities are also going to drive the demand in Bio plastics and Biopolymers market. Due to these concept there is an increase in demand of Bio plastics and Biopolymers for waste and water management and renewable energy management among others.

Increased awareness and stringent laws across the globe, has increased the demand of Bio plastics and Biopolymers.

Also frequent price fluctuation of petroleum is also contributing to the growth of the Bio plastics and Biopolymers market. Important factor is that this is recyclable, environment friendly and exhibits similar properties as petroleum based PET. Increase in demand of ecofriendly bio based polymer packaging is actually boosting up the growth of Bio plastics and Biopolymers market. Large and smooth distribution network along with strong distribution partner has led to the growth of increase in demand of Bio plastics and Biopolymers market. Rising concern for green environment and environmental safety has compelled the governments of various developed and developing nations to take safety measures which in turn has increased the usage and demand of Bio plastics and Biopolymers market significantly.

Source: LANEWS.Org - http://www.lanews.org

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March Plastics Regulations Date: March 14–15, 2018 Location: Cologne - Germany Website: www.amiplastics.com/events/event?Code=C874 Conductive Plastics Date: March 20–21, 2018 Location: Pittsburgh - USA Website: www.amiplastics.com/events/event?Code=C892 Polymers: Design, Function and Application Date: March 22–23, 2018 Location: Barcelona - Spain Website: sciforum.net/conference/polymers-2018 Plástico Brasil Date: March 25–29, 2018 Location: São Paulo - SP Website: www.plasticobrasil.com.br

April Fire Retardants in Plastics Date: April 10–11, 2018 Location: Pittsburgh - USA Website: www.amiplastics.com/events/event?Code=C0881 PVC Formulation Date: April 10–12, 2018 Location: Cologne - Germany Website: www.amiplastics.com/events/event?Code=C865 Plastic Pipes in Infrastructure Date: April 17–18, 2018 Location: London - UK Website: www.amiplastics.com/events/event?Code=C0885 PLASTEC New England Date: April 18–19, 2018 Location: Boston - USA Website: plastec-new-england.plasticstoday.com

May 5th International Conference on Plastics, Rubber and Composites (ICPRC 2018) Date: May 4–5, 2018 Location: Phuket - Thailand Website: www.icprc.org The International Plastic Showcase Date: May 7–11, 2018 Location: Orlando - USA Website: www.npe.org 34th International Conference of the Polymer Processing Society (PPS-34) Date: May 21-25, 2018 Location: Taipei - Taiwan Website: www.pps-34.com.

World Congress on Biopolymers and Polymer Chemistry Date: May 28–30, 2018 Location: Osaka - Japan Website: biopolymerscongress.conferenceseries.com

June Polymers and Organic Chemistry (POC 2018) Date: June 4–7, 2018 Location: Montpellier - France Website: iupac.org/event/polymers-organic-chemistry-2018poc-2018 4th Functional Polymeric Materials Conference Date: June 5–8, 2018 Location: Nassau - Bahamas

Website: www.fusion-conferences.com/conference76.php PLASTEC East Date: June 12-14, 2018 Location: New York – USA Website: plastec-east.plasticstoday.com Polymer Gels and Networks Date: June 17-21, 2018 Location: Prague - Czech Republic Website: www.imc.cas.cz/sympo/82pmm_png2018 Polymer Foam Date: June 19-20, 2018 Location: Pittsburgh - USA Website: www.ami.international/events/event?Code=C883 8th World Congress on Biopolymers Date: June 28-30, 2018 Location: Berlin - Germany Website: biopolymers.conferenceseries.com

July IUPAC World Polymer Congress (Macro 2018) Date: July 1-5, 2018 Location: Cairns – Australia Website: www.macro18.org

August 6th International Conference & Exhibition on Advanced & Nano Materials (ICANM 2018) Date: August 6-8, 2018 Location: Quebec - Canada Website: icanm2018.iaemm.com 3rd International Conference on Material Engineering and Smart Materials (ICMESM 2018) Date: August 11-13, 2018 Location: Okinawa - Japan Website: www.icmesm.org

September 4th International Conference on Bio-based Polymers and Composites (BiPoCo 2018) Date: September 2-6, 2018 Location: Balatonfüred - Hungary Website: bipoco2018.hu 10th Conference of Modification, Degradation and Stabilization of Polymers (MoDeSt2018) Date: September 2-6, 2018 Location: Tokyo - Japan Website: biz.knt.co.jp/tour/2018/modest/index.html Thermosetting Resins 2018 Date: September 25-27, 2018 Location: Berlin - Germany Website: thermosetting-resins.de

October 8th International Conference and Exhibition on Biopolymers and Bioplastics Date: October 15-16, 2018 Location: Las Vegas – USA Website: biopolymers-bioplastics.conferenceseries.com 8th International Conference on Polymer Science and Engineering Date: October 15-16, 2018 Location: Las Vegas – USA Website: polymerscience.conferenceseries.com

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A A A A A A A A A A A A A A A A A A A A A

ABPol Associates Sponsoring Partners

Institutions UFSCar/ Departamento de Engenharia de Materiais, SP SENAI/ Serviço Nacional de Aprendizagem Industrial Mario Amato, SP UFRN/ Universidade Federal do Rio Grande do Norte, RN

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ABPol Associates Collective Members A. Schulman Plásticos do Brasil Ltda.

Aditive Plásticos Ltda.

Avamplas – Polímeros da Amazônia Ltda.

CBE – Grupo Unigel Colorfix Itamaster Indústria de Masterbatches Ltda.

Cromex S/A Cytec Comércio de Materiais Compostos e Produtos Químicos do Brasil Ltda.

Formax Quimiplan Componentes para Calçados Ltda.

Imp. e Export. de Medidores Polimate Ltda.

Innova S/A Instituto de Aeronáutica e Espaço/AQI Jaguar Ind. e Com. de Plásticos Ltda Master Polymers Ltda.

Milliken do Brasil Comércio Ltda.

MMS-SP Indústria e Comércio de Plásticos Ltda.

Nexo International Ltda.

Nitriflex S/A Ind. e Com. Politiplastic Politi-ME. Premix Brasil Resinas Ltda.

QP - Químicos e Plásticos Ltda.

Radici Plastics Ltda.

Replas Comércio de Termoplásticos Ltda.

Uniflon - Fluoromasters Polimeros Ind .Com. Imp. Export.Ltda

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14th Brazilian Polymer Congress - 14th CBPol Águas de Lindóia - SP, October 22 to 26, 2017 Organized biannually, the Brazilian Polymer Congress is the largest event in the field of polymers. The 14th edition organized by the Engineering School of São Carlos - USP in partnership with ABPol. The congress was held at the Majestic Hotel Events Center located in the city of Águas de Lindóia, in the State of São Paulo. The city is located in the “Circuito das Águas Paulistas” region.

The opening of the 14th CBPol (Figure 1) was conducted by the ABPol president, Prof. Luiz Antonio Pessan and Dr. José Donato Ambrósio, Vice-President of ABPol; Prof. Antonio José Felix de Carvalho, General Coordinator of the 14th CBPol; Prof. Marcelo Aparecido Chinelatto, Vice-coordinator of the 14th CBPol; Prof. Cesar Liberato Petzhold, Coordinator of the Scientific Committee; Prof. Marco Aurélio De Paoli, President of the ABPol Committee for the Prize Profa. Eloisa Mano; and Prof. Edvani Curti Muniz, representing the ABPol committee for the ABPol Prize in Polymer Technology.

In the opening ceremony, as is the tradition of CBPol, the prizes ABPol Profa Eloisa Mano and ABPol of Polymer Technology were conferred, whose winners were respectively Profa. Bluma Guenther Soares from UFRJ and Eng. Paolo de Filippis from Wortex (Figure 2). After the opening ceremony followed an artistic presentation by the group “Entwined Choro Enturmado e Convidados” and then a welcome cocktail was offered to the participants (Figure 3).

Figure 1: Opening Ceremony of the 14th CBPol. From left to right: Prof. Edvani Curti Muniz, Prof. Marco Aurélio De Paoli, Prof. Antonio José Felix de Carvalho, Prof. Luiz Antonio Pessan, Dr. José Donato Ambrósio, Prof. Cesar Liberato Petzhold, and Prof. Marcelo Aparecido Chinelatto.

Figure 2: Prizes ABPol Profa. Eloisa Mano to Profa. Bluma Guenther Soares (left) and ABPol of Polymer Technology to Eng. Paolo de Filippis (right).

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Figure 3: Artistic presentation of the group “Choro Enturmado e Convidados” (left) and welcome cocktail of the 14th CBPol (right).

In this edition of CBPol, 7 plenary lectures were presented by renowned Brazilian and foreign researchers (Table 1 and Figure 4). Of the 938 papers received, 874 were approved. Of this total, 167 papers were presented as oral presentations, being 70 as keynotes presentations. 548 works were presented in the form of poster in three sessions with duration of 90 min. The evaluation of the works was coordinated by Prof. Cesar Liberato Petzhold with the collaboration of 23 area coordinators and 208 reviewers.

Thirteen companies exhibited their products and services: TA Instruments, Waters, Altmann, Anton Paar, AX Plasticos, Bruker, dpUNION, PolyAnalytik, Malvern, Scientific Mars, PerkinElmer, KCEN and Reoterm. The companies offered several technical lectures that had a great interest of the participants.

On Monday, at the end of the activities, a cocktail was offered by the organization of the event, with the objective and to promote interaction among the participants. On Tuesday, the Meeting of the Editorial Board of Polímeros Magazine: Science and Technology took place. On Wednesday, a dinner party was held for 370 members of the Pentagon hall located on the top floor of the Majestic Hotel’s Event Center.

Table 1: Plenary conferences at 14th CBPol Title

Lectures

Engineering of template microstructures derived from nanofibers Sadhan Jana – University of Akron – USA and mesoporous polymers ABPol Award “Profa. Eloisa Mano” Use of ionic liquids in Materials Bluma Güenther Soares – Federal University of Rio de Janeiro – Science Brazil Miniaturization is a trend but how does it affect the break–up and Paula Moldenaers – KU Leuven – Belgium coalescence of droplets in shear flow?

Polymers and the Smart Machines

Osvaldo Novais de Oliveira Junior – São Paulo State University – Brazil

Sustainable production of nanocellulose and their application in Julien Bras – Grenoble INP – Pagora – France fibre based packaging ABPol Award “Polymers Technology” My trajectory in the area of Eng. Paolo De Filippis - Wortex polymer processing The application of the thermoreversible furan/maleimide diels–alder Alessandro Gandini – Grenoble Institute of Technology- France/ reaction to oligomers and polymers from renewable resources EESC/USP – Brazil

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Figure 4. Plenary sessions of the 14th CBPol. From top to bottom from left to right: Prof. Sadhan Jana, Prof. Julien Bras, Eng. Paolo De Filippis, Prof. Alessandro Gandini, Prof. Osvaldo Novais de Oliveira Junior, Profa Paula Moldenaers and Profa. Bluma Güenther Soares.

The themes for submitting the congress manuscripts were defined by the organizing committee of the event that took into account the evolution of the number of works by area of previous events, as well as the current trends in the area of materials. The subjects for submission of papers were: (1) Biopolymers (cellulose / nanocelluloses, chitin / chitosan, starch, lignin and others); (2) Polymers for engineering applications; (3) Polymers for biomedical applications; (4) Polymers derived from renewable sources and biodegradable polymers; (5) Polymeric blends; (6) Rubbers and elastomers; (7) Composite Materials; (8) Polymeric nanocomposites; (9) Application of polymers in nanotechnology; (10) Recycling, degradation and stabilization of polymers; (11) Membranes and polymeric barriers; (12) Conductive polymers; (13) Rheology and processing of polymers; (14) Synthesis and modification of polymers; (15) New characterization techniques and market innovations; (16) Other areas; (17) Technical lectures. Figure 5 shows the number of work submitted for each theme. Topics 1 and 4 (biopolymers and polymers derived from renewable sources) confirm the tendency of natural polymers that was the first theme in number of works, followed by composite materials and nanocomposites. A comparison with the two previous editions of CBPol confirms the growth trend in these areas, but at the same time it does not show any drastic discontinuity in any of the themes.

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Figure 5: Number of papers submitted in the themes of the 14th CBPol. 1 Biopolymers (cellulose / nanocelluloses, chitin / chitosan, starch, lignin and others) 2 Polymers for engineering applications 3 Polymers for biomedical applications 4 Polymers derived from renewable sources and biodegradable polymers 5 Polymer Blends 6 Rubbers and elastomers 7 Composite Materials 8 Polymer nanocomposites 9 Application of polymers in nanotechnology 10 Recycling, degradation and stabilization of polymers 11 Membranes and polymer barriers 12 Conductive Polymers 13 Rheology and processing of polymers 14 Synthesis and modification of polymers 15 New Characterization Techniques and Market Innovations 16 Other areas 17 Technical lectures.

CBPol is the largest event in the field of polymers in Brazil and Latin America and receives works from all regions of Brazil and abroad as well. In order to allow a view of their regional coverage, the data of works received by each federative region are presented in Figure 6. These data are very useful, since they can serve as guides of actions to promote the development of regions that by their economic and strategic importance could have a greater participation in CBPol.

Figure 6: Number of papers presented in the 14th CBPol from the different federative regions of Brazil The works presented in poster form by students were evaluated and awarded by ABPol (first, second and third places) and by Editora Wiley (first placed in each category). This award is traditional in the CBPol and aims to value the works presented in poster format and encourage the participation of students who in general make up 60% of the CBPol congress members. In this edition the evaluation of the works was coordinated by Prof. Walter Rugerri Waldman who innovated in methodology by employing evaluation sheets and performing two job evaluations. The award-winning students are listed below by category: Polímeros, 27(4), 2017

E11

Doctorate degree 1st place - André Luís Marcomini “Obtaining and characterizing PVDF composites with high dielectric constant perovskite for supercapacitors” (UFSCar). Co-authors: Jeferson Almeida Dias, Márcio Raymundo Morelli, Rosario Elida Suman.

2nd place - Rafaela Iora Stock “Merrifield resin functionalized with optical device for the detection of cyanide” (UFSC). Co-authors: Juliana Priscila Dreyer, Vanderlei Gageiro Machado.

3rd place - Maria Verônica Silva Pinto “Resistance to impact of jute / cotton fabric composites reinforcing polyester matrix via RTM (Resin Transfer Molding)” (UFCG). Co-authors: Clarissa Coussirat Angrizani, Sandro Campos Amico, Antonio Gilson Barbosa Lima.

Master 1st place - Gean Henrique Marcatto Oliveira “Mechanical tensile properties of polypropylene composites with different cellulose structures prepared by solid state shear spray (S3P)” (UFSCar). Gean Henrique Marcatto Oliveira.

Co-authors: Leonardo Bresciani Canto, Alessandra de Almeida Lucas.

2nd place - Márcio Conti Takahashi “Compatibility of PLA composites and coconut fibers via insertion of chemically modified PLA with maleic anhydride by reactive extrusion” (UFSCar). Co-authors: Talita Rocha Rigolin, Drieli Litsue Kondo, Silvia Helena Prado Bettini.

3rd place - Pedro Vinicius de Assis Bueno “Adsorption of bovine serum albumin on polyelectrolytes and magnetic particles” (USP). Co-authors: Karina Cássia Pandeló Hilamatu, Denise Freitas Siqueira Petri.

Undergraduate sutudents 1st place - Luiz Henrique Cavalcante Damacena “Study of the interaction and migration of carvacrol and eugenol oils incorporated in the nanocomposite (LDPE / Nanoargila)” (Fatec Mauá / UFABC). Co-authors: Giovanna Amábile Duarte Rosa, Anderson Maia, Derval dos Santos Rosa, Rondes Ferreira da Silva Torin.

2nd place - Samir Leite Mathias “Characterization and production of cellulose nanocrystals of different phenological stages of the maize plant (zea mays L.)” (UFABC). Co-authors: Aparecido Junior de Menezes, Rafael Longaresi, Marcelo de Assunção Pereira da Silva.

3rd place - Matheus Fialho Zawacki “Synthesis of ATRP primers aiming to obtain styrene macromonomers” (UFRGS). Co-authors: Cesar L. Petzhold, Douglas Gamba, Luiza Cittolin Lenz, Jéssica Francielle Teixeira Chaves Petry The staff of the German Academic Exchange Service (DAAD) with the so-called “Research in Germany”, whose objective was to present the possibilities of research activity, also participated with a stand in the exhibition area and with activity at lunch time on Tuesday, with funding in Germany for students.

One aspect that marked the congress was the great participation of the congressmen in all the activities and the good level of the presented works. The social activities also contributed to a maximum interaction among the participants, providing the conditions for the establishment of partnerships and scientific interaction between congressmen and research groups based on the 14th CBPol. At the end of the congress the closing was done with a balance sheet of the congress and the General Assembly of ABPol, at which time the new board of ABPol was announced for the biennium 2018-2019.

We need to mention that the 14th CBPol was organized in the course of a strong economic crisis that passes the country and that was a constant concern throughout the preparation of the event, requiring a lot of effort to ensure that the event could be the best way possible.

We would like to thank the entire Organizing and Scientific Committee as well as the theme coordinators and the ad hoc consultants who ensure a submission process that went smoothly and most fundamentally an excellent scientific level.

We would also like to thank the agencies that supported the event, FAPESP, CAPES and CNPq, as well as the sponsors and exhibitors And finally we thank ABPol and its management for the support they have given us and Aptor for their competent work. We would also like to thank the staff of the ABPol secretariat, Marcelo Perez Gomes, Charles Fernandes de Souza and Marina Carmo Bueno for the commitment they had during the entire preparation, implementation and post‑congress, since for the organization of the congress the activities do not end at the end of the congress, but goes up to the beginning of the next, whose organization and location will be decided at the beginning of 2018.

We salute everyone who participated in the 14th CBPol and who contributed to the success of the congress and we hope to see you soon in the 15th CBPol! Antonio J. F. de Carvalho General coordinator of the 14º CBPol E12

Polímeros, 27(4), 2017

http://dx.doi.org/10.1590/0104-1428.2267

Nanostructured magnetic alginate composites for biomedical applications Pedro Marins Bedê1, Marcelo Henrique Prado da Silva1*, André Bem-Hur da Silva Figueiredo1 and Priscilla Vanessa Finotelli2 Instituto Militar de Engenharia – IME, Rio de Janeiro, RJ, Brazil 2 Faculdade de Farmácia, Centro de Ciências da Saúde – CCS, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil 1

*marceloprado@ime.eb.br

Abstract This is a study of the preparation and characterization of polymeric-magnetic nanoparticles. The nanoparticles used were magnetite (Fe3O4) and the chosen polymers were alginate and chitosan. Two types of samples were prepared: uncoated magnetic nanoparticles and magnetic nanoparticles encapsulated in polymeric matrix. The samples were analyzed by XRD, light scattering techniques, TEM, and magnetic SQUID. The XRD patterns identified magnetite (Fe3O4) as the only crystalline phase. TEM analyses showed particle sizes between 10 and 20nm for magnetite, and 15 and 30nm for the encapsulated magnetite. The values of magnetization ranged from 75 to 100emu/g for magnetite nanoparticles, and 8 to 12emu/g for coated with chitosan, at different temperatures of 20K and 300K. The saturation of both samples was in the range of 49 to 50KOe. Variations of results between the two kinds of samples were attributed to the encapsulation of magnetic nanoparticles by the polymers.

Keywords: alginate, chitosan, composite, nanoparticles, magnetic.

1. Introduction Major advances in science have allowed the nanotechnology emerge as one of the most promising research areas, in especially for their potential biomedical applications.

This paper focuses on developing a route for the production of magnetic nanoparticles encapsulated in a polymeric matrix, providing the basis from which the applied research can be stablished. Among the biomedical applications of magnetic nanoparticles, controlled drugs delivery and hyperthermia in cancer treatment are possible trends. Controlled drug delivery systems can be defined as those in which the active agent is released regardless of external factors and have well-established kinetics. These delivery systems offer several advantages when compared to other conventional systems, such as increased therapeutic efficacy, along with progressive and controlled drug release from the matrix degradation or externally controlled by diffusion, significant reduction of the toxicity and increasing remaining time in the bloody system, varied nature and composition of the vehicles and, contrary from what it might be expected, there is no predominance of instability and decomposition mechanisms of the drug (premature bio-inactivation), safe and convenient administration, with few doses and without local inflammatory reactions, directing to specific targets without significant immobilization of bioactive species, both substances hydrophilic and lipophilic can be incorporated[1].

Hyperthermia is a promising therapy for cancer treatment.

The components involved in this therapy are magnetic materials such as iron oxides (ferrofluids) and techniques applying an oscillating magnetic field. The application of this field in magnetic fluids are able to generate heat by converting magnetic energy into heat, raising the local temperature up to 41-46°C. This temperature range can

Polímeros, 27(4), 267-272, 2017

kill tumor cells without killing normal cells[2]. According to Zhao et al.[2], the magnetic particles convert the energy of the oscillating magnetic field into heat by physical mechanisms and the efficiency of this conversion is strongly dependent on the external field frequency and on the particles nature.

This heating is dependent on the size and microstructure of the particles, as these characteristics will deeply influence their magnetic properties[3]. The tumor cells are sensitive to temperature variations. Thus, in the presence of ferrofluids and applied magnetic field, the magnetic nanoparticles can eliminate tumor cells in vivo and in vitro by hyperthermia[4].

This application requires that the magnetic nanoparticles show high magnetization values for high values of thermal energy, being these particles smaller than 50 nm, with a narrow particle size distribution. Furthermore, for application in hyperthermia, these magnetic nanoparticles require special surface coating, that should be not only non-toxic and biocompatible, but also allow the targeting of particles to a specific area. Because of their hydrophobic surfaces and large surface area relative to volume, the magnetic nanoparticles tend to agglomerate and be released quickly by movement. It is possible to avoid this effect by coating the nanoparticles with biocompatible polymers[5].

Among magnetic nanoparticles, the ones that have attracted attention in biomedical applications are iron oxide nanoparticles, more precisely magnetite (Fe3O4) and maghemite (γ-Fe2O3). Magnetite ores are the most used source for obtaining iron. This ore is a mix of iron oxides, with FeO and Fe2O3 having spinel structure of inverted O2- ions with cubic packing, the larger Fe2+ ions in the octahedral interstices, half of Fe3+ ions in octahedral sites and half of

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O O O O O O O O O O O O O O O O

Bedê, P. M., Silva, M. H. P., Figueiredo, A. B.-H. S. & Finotelli, P. V. the remaining in tetrahedral positions. The magnetization of Fe3O4 occurs in the presence of external magnetic field, and disappears when the field is removed. This effect is due to non-conservation of magnetic orientation of individual atoms[6]. As for the polymeric matrix, the alginate is a very consistent choice, because it is biocompatible, allows its use in biomedical applications. It also forms a strong gel in the presence of divalent cations, especially calcium, by ionic crosslinks between the polyionic alginate chains[7].

This gelation of alginate is conventionally described in terms of the “egg box” model, where divalent cations are coordinately bound to the carboxyl groups of guluronic acid[8]. This model is shown in Figure 1. The structure of guluronic acid provides the right distance between the carboxyl and hydroxyl groups, with a high coordination degree with calcium Ca cations[10].

FeCl2 and FeCl3 solutions were mixed and kept under mechanical stirring at 60°C during 15 minutes. The next step was to drop 200 mL of 25% ammonium hydroxide (NH4OH) were drop into the FeCl2 + FeCl3 mixture. The precipitate was washed with 50 ml of distilled water, centrifuged and lyophilized.

2.1 Encapsulation of FE3O4 in the alginate/chitosan matrix The first step was to prepare three different solutions: - Sodium alginate – 9.5 ml, 0.06%; - Calcium chloride – 0.5 ml, 18 mM; - Chitosan – 2 ml, 0.05%.

2. Materials and Methods

Sodium alginate solution was mixed with 0.0024 g of magnetite.

The materials used in all stages of polymer-magnetic nanoparticles were purchased from Fluka – Biochemika (sodium alginate), Aldrich (chitosan), and VETEC (calcium chloride and ammonium hydroxide).

Magnetite was obtained by the co-precipitation method from Fe II and Fe III in an alkaline medium according to the reaction:

The alginate + Fe3O4 solution was constantly sonicated (probe sonicator), so that complete mixing of the solutions could occur. CaCl2 solution was slowly dripped and sonicated.

After complete mixing, chitosan was added to the calcium alginate and magnetite mixture, and then sonicated, remaining under stirring during further 20 min and centrifuged at 6000 rpm at 24 °C during 30 min. The scheme of this process can be seen in Figure 2.

2FeCl3 + FeCl2 + 8NH 4Cl → 8NH 4OH + Fe3O 4 + 4H 2O (1)

In order to characterize nanoparticles, the following techniques were applied: X-Ray Diffraction was performed on X’Pert PRO (Panalytical) equipment. The method used was the powder method, using copper Kα1 radiation with wavelength (λ) 1.54056 Å. The measurements were performed with 40 mA and 40 kV. Size and Zeta Potential was performed on a ZETA PLUS ANALYZER equipment, from BROOKHAVEN INSTRUMENTS CORPORATION company. The ZETA PLUS parameters used were: five runs of 30 seconds each in ultrapure water with refractive index 1.340.

Figure 1. “Egg-box” model for Calcium Alginate[9].

Transmission Electron Microscopy was performed on a FEI Tecnai G20 equipment, with 200 kV voltage.

Figure 2. Schematic illustration of the magnetite encapsulation process.

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Polímeros, 27(4), 267-272, 2017

Nanostructured magnetic alginate composites for biomedical applications Magnetic Measurements (SQUID) was performed using Quantum Design MPMS-5S SQUID magnetometer.

3. Results and Discussion 3.1 X-Ray Diffraction (XRD) X-Ray diffraction analyses were performed for Fe3O4 and Fe3O4 encapsulated in alginate/chitosan complex polymer. Figure 3 shows the diffraction pattern characteristic of magnetite (JCPDS 19-0629). Figure 4 presents the diffraction pattern obtained for the encapsulated magnetite, where the presence of amorphous material is evident. The amorphous phase was associated to the complex alginate with chitosan.

The good definition of the reflection peaks in the Fe3O4 sample, resulting in a well-defined diffraction pattern, confirms the success of the synthesis method for the production of well crystallized magnetite as fabricated. The crystallite size was 10 nm, determinated by Scherrer’s method.

3.2 Size and zeta potential The size measurements are important to define the best method of sample preparation, since it enables the choice for a size range that will provide the desired material properties. The mean diameter of magnetic nanoparticles, as shown in Table 1, was 155.8 nm with a polydispersity average of 0.213. The diameter for the the magnetic nanoparticles encapsulated, shown in Table 2, was 255.0 nm with a polydispersity average of 0.330. These tables show a polydispersity narrow size distribution and homogeneity of the nanoparticle, indicating stability and control of the

diameter[11]. In a previous study, Ma et al.[12] synthesized nanoparticles of magnetite (Fe3O4) coated with alginate, and found an average diameter of 193.8 nm and polydispersity index of 0.209. Ahmad et al.[13] synthesized nanoparticles of alginate/chitosan and obtained with an average diameter of 229 nm polydispersity of 0.44 of these nanoparticles. It is important to emphasize that this is probably related to the diameter of agglomerates and not the isolated nanoparticles.

A possible confirmation that may arise through the analysis of transmission electron microscopy is the fact that the encapsulated nanoparticles showed higher average size than non-encapsulated may be associated with the own polymer coating these nanoparticles. Due to these small sizes, the material is expected to present the desired magnetic properties, but this can only be confirmed through the characterization of these properties.

The zeta potential measurements were performed in order to find out if the material exhibited good stability when in suspension. The general rule for electrostatic stability of the solution is the Zeta potential range of +/- 30 mV. As the zeta potential of the three samples is outside this range, we can consider that the samples are stable. An important factor to be considered is the presence of chitosan in the material.

One can perceive a potential variation between samples with and without chitosan, the ones with chitosan being more stable. This can be a direct consequence of the fact that chitosan acts as an inhibitor of surface charges existing in the alginate, thereby contributing to a better stability.

Table 3 shows the change in nanoparticle Zeta potential between the coated and non-coated samples.

Table 1. Size of magnetic NP. MEASUREMENTS DIAMETER (nm) POLYDISPERSITY 1 150.4 0.195 2 160.9 0.242 3 145.1 0.190 4 163.4 0.209 5 159.2 0.231 AVERAGE 155.8 0.213 ERROR 3.5 0.010

Figure 3. XRD pattern of magnetic nanoparticles.

Table 2. Size of magnetic NP encapsulated in polymeric matrix.

MEASUREMENTS DIAMETER (nm) POLYDISPERSITY 1 230.0 0.677 2 248.7 0.199 3 260.6 0.327 4 271.9 0.260 5 263.7 0.184 AVERAGE 254.9 0.329 ERROR 7.3 0.091

Table 3. Zeta potential of the nanoparticles.

Figure 4. XRD pattern of magnetic nanoparticles encapsulated in polymeric matrix.

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pH 3.61 3.61 3.61

Sample PM-NP w/ chitosan PM-NP Magnetic NP

Average Zeta Potential -(37.59 ± 2.9) mV -(30.63 ± 0.8) mV -(30.45 ± 2.7) mV

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Bedê, P. M., Silva, M. H. P., Figueiredo, A. B.-H. S. & Finotelli, P. V. 3.3 Transmission Electron Microscopy (TEM) Figure 5 refers to magnetic nanoparticles without encapsulation. A nearly spherical shape and a strong trend to agglomerate can be observed. The particles diameter is between 10 and 20 nm, which allows these nanoparticles to present superparamagnetic behavior. In a previous work, Kim et al.[5] synthesized Fe3O4 nanoparticles coated with chitosan and observed average diameter of 10.3 nm.

Figure 6 refers the magnetic nanoparticles encapsulated in polymeric matrix (PM-NP). It can be observed that the coated particles were more dispersed, when compared

to uncoated ones. This finding may be linked to the fact that magnetic nanoparticles are encapsulated in a matrix of alginate and chitosan; chitosan, acting as an inhibitor of surface charges, has resulted in a better stability of the nanoparticles. Regarding the size, the observed diameter was in the range of 15 to 30nm, which allows these PN-PM to present superparamagnetic behaviors. TEM analyses also confirmed that the diameter size found in the analysis by light scattering techniques is related to the clusters of nanoparticles. In a previous study, Ma et al.[12] observed average diameter of 10nm in NP-PM magnetite (Fe3O4) coated with alginate. Energy dispersive spectroscopy (EDS) was also carried out for the PM-NP nanoparticles, highlighted by the red rectangle region in Figure 6. The presence of the elements Ca, Fe and C was confirmed.

3.4 Magnetic measurements (SQUID) The magnetization measurements of samples A and B were performed at temperatures 20 K and 300 K. Figure 7 refers to uncoated samples, while Figure 8 refers to coated samples. At 20 K, the saturation field of the samples A and B was approximately 50 KOe, and 300 K of 49 KOe. The results for the maximum magnetization of the Fe3O4 nanoparticles

Figure 5. Electromicrograph (TEM) of uncoated magnetic nanoparticles.

Figure 6. Electromicrograph (TEM) of magnetic nanoparticles encapsulated in polymeric matrix (PM-NP). The rectangle shows the area from which EDS was performed.

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Figure 7. MxH curve at 20K (a) and 300K (b) of magnetite nanoparticles.

Polímeros, 27(4), 267-272, 2017

Nanostructured magnetic alginate composites for biomedical applications

4. Conclusions With the intention to produce encapsulated magnetic particles for possible use biomedical applications, alginate and magnetite nanoparticles were successfully synthesized.

The nature of the iron oxide (magnetite, Fe3O4) was confirmed by identification of diffraction peaks in XRD analysis.

Crystallite size was determined by the Debye-Scherrer method and found an average diameter of 10nm. The material presented good stability, measured by the zeta potential and ideal size, which could be measured by light scattering and transmission electron microscopy. Microscopy also allowed to provide a morphological analysis of the material, confirming the expected nearly spherical shape of the nanoparticles.

The good magnetization and superparamagnetic character of the material was confirmed by the measurements of the magnetic SQUID.

5. Acknowledgements The authors would like to acknowledge to CNPq, CAPES, CEPEL and CBPF.

6. References

Figure 8. MxH curve at 20K (a) and 300K (b) of encapsulated magnetite nanoparticles.

without coating ranged from 75 to 100 emu/g at 300 K and 20 K, respectively, and Fe3O4 nanoparticles in the coating ranged from 8 to 12 emu/g at 300 K and 20 K, respectively.

A decrease in the magnetization was observed at 300 K for both samples. A further reduction was also observed, when comparing coated and uncoated samples. This reduction may be explained by the encapsulation, being the polymers a barrier to magnetization. If the alginate completely covers Fe3O4 nanoparticles, the magnetization will drop significantly[14].

This reduction can also be associated with two factors: the drying of the encapsulated samples, and especially the concentration of Fe3O4 used in their preparation. Another important fact to note is that both samples at 20 K and 300 K, showed virtually no hysteresis, presenting coersivity almost nil. This may confirm the superparamagnetic character of the samples. In a previous study, Ma et al.[12] found values for the magnetization of magnetite nanoparticles (Fe3O4) encapsulated in alginate matrix ranging from 30 to 55 emu/g at room temperature. Ma et al.[12] associated with this variation to the concentration of Fe3O4 used in the synthesis of their samples; the saturation field found by Ma et al.[12] in all samples was slightly higher than 10 KOe, a value almost six times smaller when compared to those found in the present study.

Polímeros, 27(4), 267-272, 2017

1. Kumar, R. (2000). Nano and microparticles as controlled drug delivery devices. Journal of Pharmacy & Pharmaceutical Sciences, 3(2), 234-258. PMid:10994037.

2. Zhao, D., Zeng, X., Xia, Q., & Tang, J. (2006). Inductive heat property of Fe,O, nanoparticles in AC magnetic field for local hyperther mia. Rare Metals, 25(6), 621-625. http://dx.doi. org/10.1016/S1001-0521(07)60159-4.

3. Atsumi, T., Jeyadevanb, B., Satob, Y., & Tohji, K. (2007).

Heating efficiency of magnetite particles exposed to AC magnetic field. Journal of Magnetism and Magnetic Materials, 310(2), 2841-2843. http://dx.doi.org/10.1016/j.jmmm.2006.11.063.

4. Jordan, A., Scholz, R., Wust, P., Fähling, H., & Roland Felix (1999). Magnetic fluid hyperthermia: cancer treatment with AC magnetic field induced excitation of biocompatible superparamagnetic nanoparticles. Journal of Magnetism and Magnetic Materials, 201(1-3), 413-419. http://dx.doi. org/10.1016/S0304-8853(99)00088-8.

5. Kim, D. H., Lee, S. H., Im, K. H., Kim, K. N., Kim, K. M., Shim, I. B., Lee, M. H., & Lee, Y.-K. (2006). Surface-modified magnetite nanoparticles for hyperthermia: preparation, characterization, and cytotoxicity studies. Current Applied Physics, 6(S1), 242246. http://dx.doi.org/10.1016/j.cap.2006.01.048.

6. Sidhu, P. S., Gilkes, R. J., & Posner, A. M. (1978). The synthesis and some properties of Co, Ni, Zn, Cu, Mn and Cd substituted magnetites. Journal of Inorganic and Nuclear Chemistry, 40(3), 429-435. http://dx.doi.org/10.1016/0022-1902(78)80418-7.

7. Boisseson, M. R., Leonard, M., Hubert, P., Marchal, P., Stequeart, A., Castel, C., Favre, E., & Dellacherie, E. (2004).

Physical alginate hydrogels based on hydrophobic or dual hydrophobic/ionic interactions: Bead formation, structure, and stability. Journal of Colloid and Interface Science, 273(1), 131-139. PMid:15051442. http://dx.doi.org/10.1016/j. jcis.2003.12.064.

8. Iskakov, R. M., Kikuchi, A., & Okano, T. (2002). Timeprogrammed pulsatile release of dextran from calcium-alginate gel beads coated with carboxy-n-propylacrylamide copolymers.

Journal of Controlled Release, 80(1-3), 57-68. PMid:11943387. http://dx.doi.org/10.1016/S0168-3659(01)00551-X. 9. Shao, F., Ankur, T., Diana, M. S., Riccardo, L. B., Ira, S. B., Sachin, V., Eric, J. M., & Lawrence, H. B. (2011). Relevance 271

Bedê, P. M., Silva, M. H. P., Figueiredo, A. B.-H. S. & Finotelli, P. V. of rheological properties of sodium alginate in solution to calcium alginate gel properties. AAPS PharmSciTech, 12(2), 453-460. PMid:21437788. http://dx.doi.org/10.1208/s12249011-9587-0.

10. Finotelli, P. V. (2006). Microcápsulas de alginato contendo nanopartículas magnéticas para liberação controlada de insulina (Tese de doutorado). Instituto de Química, Universidade Federal do Rio de Janeiro, Rio de Janeiro.

11. Kulkamp, I. C., Paese, K., Guterres, S. S., & Pohlmann, A. R. (2009). Estabilização do ácido lipoico via encapsulação em nanocápsulas poliméricas planejadas para aplicação cutânea.

Quimica Nova, 32(8), 2078-2084. http://dx.doi.org/10.1590/ S0100-40422009000800018.

12. Ma, H., Qi, X. R., Maitani, Y., & Nagai, T. (2007). Preparation and characterization of superparamagnetic iron oxide nanoparticles stabilized by alginate. International Journal of

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Pharmaceutics, 333(1-2), 177-186. PMid:17074454. http:// dx.doi.org/10.1016/j.ijpharm.2006.10.006.

13. Ahmad, Z., Pandey, R., Sharma, S., & Khuller, G. K. (2006).

Pharmacokinetic and pharmacodynamic behaviour of antitubercular drugs encapsulated in alginate nanoparticles at two doses. International Journal of Antimicrobial Agents, 27(5), 409-416. PMid:16624533. http://dx.doi.org/10.1016/j. ijantimicag.2005.12.009.

14. Denizot, B., Tanguy, G., Hindre, F., Rump, E., & Jeune, J. J. L. & Jallet, P. (1999). Phosphorylcholine coating of iron oxide nanoparticles. Journal of Colloid and Interface Science, 209(1), 66-71. PMid:9878137. http://dx.doi.org/10.1006/ jcis.1998.5850.

Received: Aug. 11, 2015 Revised: Mar. 11, 2016 Accepted: May 17, 2016

Polímeros, 27(4), 267-272, 2017

http://dx.doi.org/10.1590/0104-1428.2397

Green polyurethane synthesis by emulsion technique: a magnetic composite for oil spill removal Raphael Maria Dias da Costa1, Gabriela Hungerbühler1, Thiago Saraiva1, Gabriel De Jong1, Rafael Silva Moraes1, Evandro Gonçalves Furtado2, Fabrício Machado Silva3, Geiza Esperandio de Oliveira4, Luciana Spinelli Ferreira1,5 and Fernando Gomes de Souza Junior1,4* Laboratório de Biopolímeros e Sensores – LaBioS, Instituto de Macromoléculas – IMA, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil 2 Alfa Rio Química Ltda., Duque de Caxias, RJ, Brazil 3 Instituto de Química – IQ, Universidade de Brasília – UnB, Brasília, DF, Brazil 4 Programa de Engenharia Civil – PEC, Instituto Alberto Luiz Coimbra de Pós-graduação e Pesquisa de Engenharia – COPPE, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil 5 Programa de Engenharia da Nanotecnologia – PENt, Instituto Alberto Luiz Coimbra de Pós-graduação e Pesquisa em Engenharia – COPPE, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil 1

*fgsj@ufrj.br

Abstract After the consolidation of the Brazilian biodiesel industry, issues related to the final destination of the glycerin, the by-product from the biodiesel industrial process, drawing the attention of several researchers. There are several uses to this byproduct. Among them, the obtaining of polymers, such as polyurethane (PU), are very encouraged since the glycerin ca be used, as well as the castor oil, in the replacement of petrochemical polyols. The aim of this work was to propose a new route for the obtainment of a petroleum sorbent based on polyurethane resin from glycerin and castor oil, through the emulsion technique. In addition, maghemite (γ-Fe2O3) was mixed to the polymer matrix, producing a magnetic composite, able to make easier the oil cleanup process. The products synthesized were characterized by Fourier transform infrared spectroscopy, X-ray diffraction, simultaneous Thermogravimetry (TGA) and Differential scanning calorimetry (DSC), Optical microscopy, Scanning electron microscopy (SEM). In addition, magnetic force and oil removal capability tests were also performed. The magnetic material was used to remove oil from water, exhibited a good oil removal capability. In a typical test, 1g of the composite containing 5wt% of maghemite was able to remove 10g of oil from water.

Keywords: green polyurethanes, glycerin, castor oil, maghemite, magnetic composites.

1. Introduction The biodiesel consists of an alkylester of fatty acids, obtained by the transesterification of vegetable oils or animal fats, using a small chain alcohol, such as methanol or ethanol. Aiming the replacement of part of the fossil fuels employed as an energy source, the production of Brazilian biodiesel as alternative fuel has been increased during the last years[1]. The Brazilian law nº 11.097 / 2005 determined the increasingly addition of biodiesel to diesel. From the initial amount of 2%, nowadays 7% of biodiesel is added to the diesel[2]. So, the growing using of the biodiesel is responsible for the surplus production of glycerol (each 100 kg of the biodiesel can generate 10 kg of glycerol). Therefore, as long as the usage of biodiesel as alternative fuel has been encouraged by Brazilian Government, innovative uses for the glycerin source must be pursued[3].

In this context, the glycerin can be employed for the production of resins, such as polyurethanes. The polyurethanes (PU) are one of the most important polymeric materials, showing different characteristics and applications[4-6].

Among the variety of applications, PUs can be employed as oil spill cleanup agents from aqueous systems[7-9], which is a major concern since petroleum use always involves a

Polímeros, 27(4), 273-279, 2017

considerable risk of oil spillage[10,11] that could occur during its transportation and storage[12]. Thus, in a previous work from our research group, magnetic foams were prepared with the insertion of maghemite nanoparticles into a polymer matrix, aiming to make the removal of the polymer material containing petroleum from the water surface easier[9]. These foams were prepared from renewable feedstock, as castor oil and glycerin (byproduct from the biodiesel industry).

However, this material was obtained by a bulk polymerization process, in which there are relevant operational issues.

Among them, the use of higher temperatures and the need of additional steps, such as the milling of the obtained polymer, make the final process less competitive from the industrial point of view.

Therefore, the main goal of this work was the preparation of a magnetic oil sorbent, by emulsion polymerization. Presented material is a polyurethane resin obtained from renewable feedstock in the presence of maghemite nanoparticles[13].

Emulsion technique allows to perform the reaction in lower temperatures and reduced viscosity of the reaction medium in comparison to the bulk technique used previously[9].

In addition, the composite is obtained as very small particles

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O O O O O O O O O O O O O O O O

Costa, R. M. D., Hungerbühler, G., Saraiva, T., De Jong, G., Moraes, R. S., Furtado, E. G., Silva, F. M., Oliveira, G. E., Ferreira, L. S., & Souza, F. G., Jr. at the end of polymerization. The material was used in oil spill cleanup tests and results shown that each gram of PU is able to remove around ten grams of oil. This result is very encouraging since each gram of the previous material, milled after the by bulk polymerization, was able to remove four grams of the petroleum from the water.

2. Experimental Part 2.1 Materials Glycerin and sodium lauryl ether sulfate (C12H25NaO4S) were purchased from Vetec (Rio de Janeiro, Brazil).

Dodecylbenzenesulfonic acid (DBSA) was purchased from Solquim LTDA (Brazil). Castor oil was kindly donated by Petrobras, with density in the range of 0.954-0.965 (@ 25 °C).

Toluene diisocyanate (TDI, C9H6O2N2) was purchased from Bayer (Rio de Janeiro, Brazil). These reagents were used as received, without further purification.

2.2 Preparation of the emulsion Before the polymerization, the emulsion was prepared following the methodology described in the literature[14].

In a typical procedure, a solution containing 0.5 g of DBSA into 1 liter of distilled water was previously prepared, and employed as continuous phase. Castor oil, glycerin and sodium lauryl ether sulfate (SLES) were dispersed in the aqueous solution, by using a homogenizer Polytron (PT 3100 D), at 10,000 rpm for 10 minutes, using different amounts of monomers, surfactants and catalyst, which are described in the section 2.4.

2.3 Synthesis of maghemite As described in previous works[9,15-25], our group is interested in the maghemite synthesis and we used similar approaches to obtain this nano material. Specifically in the present work, aqueous solutions of hydrochloric acid (2 M), ferric chloride (2 M), and sodium sulfite (1 M) were prepared. Into a beaker, under continuous agitation, 30 mL of the ferric chloride solution and 30 mL of deionized water were added. Soon afterwards, 20 mL of the sodium sulfite solution was added to the beaker, still under continuous agitation. The reaction product was precipitated by slowly adding 51 mL of concentrated ammonium hydroxide into the beaker under continuous agitation. The medium was poured after 30 min. The obtained nanoparticles were washed several times using distilled water and finally dried to constant weight at 60 °C in an oven. After that, sample was heating at 250 °C for 1 h.

2.4 Synthesis of polymer and composite Soon afterwards the emulsion preparation, the polymerization took place at 70 °C using TDI (0.309 mol), glycerin (0.204 mol), castor oil (0.002 mol) and surfactant (0.0069 mol). Synthesis was carried out under continuous and stirring by a mechanical stirrer Fisatom (model 713D), at 800 rpm; and heated by a magnetic stirrer with heating from IKA (model C-MAG HS7) with a thermocouple attached.

During stirring, TDI was added into the reaction medium.

The product was precipitated in an excess of ethanol, and 274

filtered. The composite was prepared following the same steps, except for the inclusion of 5wt% of maghemite, which was performed by the dispersion of the filler in the medium containing the emulsion also under mechanical stirring.

2.5 Characterization Polymer, maghemite and composite were analyzed in a Fourier transform infrared with the attenuation reflectance accessory (FTIR-ATR), from Nicolet model iN10. The used range was from 4000 to 600 cm-1, with 20 scans accumulated in a resolution of 4 cm-1. Maghemite and composite were characterized by X-ray diffraction (XRD), on a Rigaku‑ULTIMA IV X-ray diffractometer, in the angular range 2θ from 2° to 80°, using CuKα radiation (k= 1.5418 Å, 30 kV and 15mA). Optical microscopy was performed in a Bel Photonics microscope, aiming the calculation of the particles diameter. Magnetic force test of maghemite/composite was performed following the procedure established in our laboratory[22]. Simultaneous TGA-DSC analysis were performed using a PerkinElmer STA-6000 in the temperature range between 28-670 °C in 20mL/min flow of nitrogen and heating rate of 10 °C/min.

The Scanning Electron Microscopy was performed using a JEOL JSM-5610 LV with acceleration voltage of 15 kV. The cure degree was performed using a sample, previously weighed, inside a Soxhlet extractor in a filter paper package.

The used solvent during the extraction process was an equivolumetric mixture of toluene, heptane and cyclohexane at 250 °C. Then, the samples were dried and weighed again to determine the weight lost and, consequently, the cure degree. For the oil removal test, in 300 mL beaker, 2 g of petroleum were spilt on a 150 mL saline water solution (55000 ppm, containing NaCl and CaCl2, in a 10:1 rate).

Then, a known amount of the magnetic resin was poured into the oily system. After 5 minutes, the oil and composite were magnetically removed, and the mass of the oil residue could be determined by gravimetry[26].

3. Results and Discussion Figure 1 shows the samples FTIR spectra of the resin, maghemite and composite filled with 5wt% of maghemite.

The FTIR spectrum of the resin is presented in the Figure 1a.

Among its characteristic bands it is possible to see a wide one related to O-H stretching of the hydroxyl group, which takes place around 3300 cm-1. The profile of this band is associated with the hydrogen bond that these compounds are able to form. The intense characteristic peak observed at 2272 cm-1 is related to the stretching of the bands N=C=O from isocyanate groups. The characteristic band located around 1650 cm-1 is C=O axial deformation from carbonyl group. While the vibration of C-N band takes place at 1540 cm-1. The presence of the remaining cyanate group indicates reactive sites which can allow further modifications in the resin surface. In the FTIR spectrum of the maghemite, presented in the Figure 1b, the main characteristic bands appears around (i) 3450cm-1, associated with the Fe-OH and O-H conjugated bands, and at (ii) 634 and 584 cm-1, related to stretching of the band Fe-O, from the phases ε-Fe2O3 e α-Fe2O3, respectively[27].

Polímeros, 27(4), 273-279, 2017

Green polyurethane synthesis by emulsion technique: a magnetic composite for oil spill removal In turn, the spectra of the composite did not present any significant difference (Figure 1c), when compared to the matrix, which can indicate the absence of significant chemical interactions between the matrix and the filler, due to the small amount of the latter.

Pure resin presents only a broad hump around 23°, which is related to the amorphous nature of this macromolecular material. The X-ray diffratogram of pure maghemite presents peaks at 30.29°, 35.39°, 42.91°, 52.97°, 57.36º and 62.86°, corresponding to (220), (311), (440), (400), (511) and (440) crystalline plane from an orthorhombic crystal system.

Crystallite size was calculated in accordance with Scherer’s Equation 1, using the sharpest peak placed at 2θ equal to 35.61°, related to the (331) plane of the maghemite (Figure 2).

Cs = K λ / ( β cos θ ) (1)

In Equation 1, Cs is the crystallite size, λ the wavelength, the full width at half maximum and θ is the Bragg angle (2θ/2).

Figure 1. FTIR spectra of PU resin (a), maghemite (b) and composite filled with 5wt% of the maghemite (c).

Figure 2. XRD of PU resin (a), maghemite (b) and composite filled with 5wt% of the maghemite (c).

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From this calculation, the crystallite size obtained for pure maghemite was equal to 6.5 ± 0.5 nm, while in the composite, its value was equal to 7.0 ± 0.6, which guarantees the nanometric size of maghemite particles inside the composite without the formation of agglomerated material.

Figure 3 shows the optical micrography and the probability density function of resin and composite particles.

Obtained results, with 95% of confidence, allow inferring that the diameters of the polymer particles ranges between 5.20 and 21.02 µm. In addition, the most probable observed value is equal to 14.06 µm. In turn, the diameters of the composite particles ranges between 3.84 and 23.31µm and the most probable observed value is equal to 13.70 µm.

The size of the samples, in both cases, was equal to 43 counts. Obtained results allow us to infer that the size of the particles remains statistically the same after the preparation of the composite material.

The results from the magnetic force test under the influence of different magnetic fields are shown in the Figure 4.

There is an increase in the magnetic force when the magnetic field is increased, proving the presence of magnetic properties. In addition, magnetic force at maximum magnetic field (equal to 780 Gauss) of the pure maghemite and composite was equal to 824.8 ± 0.02 mN/g and 735.9 ± 0.001 mN/g, respectively. The value for the composite is slightly lower than the one for the pure maghemite, due to the small amount of the filler in the polymer matrix (equal to 5wt%).

Figure 5 shows the simultaneous TGA-DSC analysis of the tested materials. Composite presents a small weight loss before 100 °C, which is related to the residual water.

The magnetic composite remains thermally stable before 227 °C, while the petroleum and the resin impregnated of petroleum are stable before 133 °C and 224 °C, respectively.

From 135 °C to 515 °C, petroleum lost 95% of its weight, thus this petroleum is mainly compounded by heavy fractions[28].

In the 280-380 °C range petroleum presented the smallest weight loss, equal to 25.6wt% while composite impregnated of petroleum and composite presented 39.5wt% and 55.3wt%, respectively. In spite of presented weight loss, petroleum did not present an increase of the weight loss rate in this range while composite and composite impregnated of petroleum did. Something similar took place in DSC results. Thus, the observed enthalpic phenomenon is mainly related to the decomposition of composite and the associated enthalpy allow us to infer that along decomposition, the studied composite absorbed 346J/g while composite impregnated of petroleum only 168J/g. This difference is due to the crude oil presence, which did not absorb heat in this temperature range. On the other hand, the second weight loss event, from 410 to 530 °C, is mainly related to the considerable presence of lubricating (C26-C38 and 4-10 rings per molecule[29]) and vacuum gas oil (mainly composed by high paraffins[30]) inside tested petroleum. TGA results between 410 and 530 °C allowed inferring that composite, composite impregnated of petroleum and petroleum presented weight loss equal to 6.8wt%, 22.9wt% and 41.8wt%, respectively. In turn, DSC results showed that petroleum and composite impregnated of petroleum presented considerable endothermic events.

The enthalpy of composite, composite impregnated of petroleum and petroleum, in the same temperature range, 275

Costa, R. M. D., Hungerbühler, G., Saraiva, T., De Jong, G., Moraes, R. S., Furtado, E. G., Silva, F. M., Oliveira, G. E., Ferreira, L. S., & Souza, F. G., Jr.

Figure 3. Optical microscopy and probability density function of the diameter of the resin (a & b) and composite (c & d).

Figure 4. Magnetic force test of the pure maghemite (a) and the composite containing 5wt% of the maghemite (b).

Figure 5. Simultaneous TGA-DSC analysis of petroleum (a & a’) composite impregnated of petroleum (b & b’) and composite (c & c’).

were equal to 9.3J/g, 21.3J/g and 84.2J/g, respectively. Thus, considering that enthalpy of composite and petroleum are equal to 0% and 100% of petroleum in the samples, the amount of remain heavy fractions on the surface of the composite can be estimated by interpolation. Therefore, calculations lead to the conclusion that there was at least 16% of petroleum absorbed on the surface of the composite. This result is in agreement with the SEM one, showed in Figure 6, which

allowed to see that studied material presented a porous surface, able to absorb the petroleum, making easier the oil removal from the water. Figure 6 also showed the SEM of the pure maghemite. The particles are agglomerated due to the drying process, which is indispensable to the obtaining of the micrography in low vacuum. In spite this agglomeration, particles presented an average diameter, calculated using the imageJ® software, equal to (108 ± 17) nm. This is an

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Green polyurethane synthesis by emulsion technique: a magnetic composite for oil spill removal

Figure 6. Scanning electron microscope image of the maghemite (a) and composite (b).

Figure 7. Oil removal from water using a magnetic composite: water with oil spilled (a), oil after the addition of the hybrid material (b) and oil and resin removed using a magnet (c).

interesting value, since allows the easy dispersion of these small magnetic particles within the emulsion phase.

The insertion of maghemite into the polymer matrix aims to make easier the cleanup process, removing both petroleum and resin from the aqueous system. The oil removal capability tests, showed in Figure 7, were performed according to the analytical procedure established in our laboratory[18].

Before oil removal tests, the prepared resin was dispersed in water, where it was kept for five minutes. After this time, the resin was collected using a nylon sieve. Obtained results showed that 1.00 g of the resin absorbs only 0.25 g of water, demonstrating the hydrophobic property of the material.

Oil removal tests using the presented magnetic composite, prepared here by emulsion, showed that each gram of the composite is able to remove (10.0 ± 0.2)g of petroleum from the water. For comparative purpose, in a previous work[9] our group produced a magnetic PU based on castor oil, toluene diisocyanate and water by a bulk polymerization.

Each gram of this material was able to remove (4.1 ± 0.1)g of petroleum from the water. Other example is the peat, a common raw material, which is able to sorb around 3.9g of petroleum when exposed to this contaminant for 5 min[31], however, without the magnetic removal capability, which speeds up the oil spill cleanup process. Thus, the material presented here has larger oil removal capability, being cheaper (due to the use of glycerin) and easier to prepare than the ours previous one. In addition, tested materials Polímeros, 27(4), 273-279, 2017

presented an average cure degree equal to (99.3 ± 0.2)%, which means that, after an oil extraction by solvents, the sorbers can be reused, as proved by three additional tests, which showed that oil removal capability remained the same. The advantages of the preparation by emulsion are very encouraging, since the material can be easily prepared and it is obtained in the granular form, avoiding the milling process. Therefore, presented material is promising to the cleanup process and environmental recovery.

4. Conclusions Previous experiments have shown the viability of obtaining polyurethane from castor oil and glycerin through an emulsion technique. The employment of alternative feedstock is a route for the synthesis of green resins.

The emulsion polymerization was able to produce crude oil absorbers, with a high cure degree. This is a technological advantage, since it can promote the reusing of the absorber material, as well as the crude oil recovery. The addition of magnetic nanoparticles into the polymer matrix allowed the production of a magnetic composite, which possesses a good magnetic force. Obtained magnetic force constitutes other technological advantage, since the use of magnetic tools make easier the cleanup process. The oil removal capability of the material is equal to (10.0 ± 0.2)g of petroleum per gram of material. This is an improvement in 277

Costa, R. M. D., Hungerbühler, G., Saraiva, T., De Jong, G., Moraes, R. S., Furtado, E. G., Silva, F. M., Oliveira, G. E., Ferreira, L. S., & Souza, F. G., Jr. the oil removal capability of this material, which allows its use in oil spill cleanup processes. Therefore, the prepared material contributes to the environment encouraging nobler uses to some available renewable resources.

5. Acknowledgements The authors thank to Conselho Nacional de Desenvolvimento Científico e Tecnológico (CNPq‑474940/2012-8 and 550030/2013-1), Coordenação de Aperfeiçoamento de Pessoal de Nível Superior (CAPES and CAPES-NANOBIOTEC), Financiadora de Estudos e Projetos (FINEP PRESAL Ref.1889/10) and Fundação Carlos Chagas Filho de Amparo à Pesquisa do Estado do Rio de Janeiro (FAPERJ) for the financial support and scholarships.

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31. Klavins, M., Porshnov, D., Ansone, L., Robalds, A., & Dreijalte, L. (2012). Peat as natural and industrial sorbent. In R. A. R. Ramos, I. Straupe & T. Panagopoulos (Eds.). Recent researches in environment, energy systems & sustainability (pp. 146-151).

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Received: Nov. 09, 2015 Revised: Apr. 12, 2016 Accepted: May 16, 2016

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http://dx.doi.org/10.1590/0104-1428.2423

O O O O O O O O O O O O O O O O

Thermal degradation of polymer systems having liquid crystalline oligoester segment Renato Matroniani1* and Shu Hui Wang1 Departament of Metallurgical and Materials Engineering, Polytechnic School, Universidade de São Paulo – USP, São Paulo, SP, Brazil

1

*rmatroniani@usp.br

Abstract Block copolymers and blends comprised by liquid crystalline oligoester and polystyrene were prepared and their thermal stability were characterized by thermogravimetric analysis (TGA). The samples have shown three main decomposition temperatures due to (1) lost of flexible chain and decomposition of mesogenic segment, (2) decomposition of polystyrene and (3) final decomposition of oligoester rigid segment. Both copolymers and polymer blends presented lower thermal stability compared to polystyrene and oligoester. The residual mass after heating at 600 °C in copolymers and polymer blends were lower than those found in the oligoesters. A degradative process of aromatic segments of oligoester induced by decomposition of polystyrene is suggested.

Keywords: blends, liquid crystalline oligoester, thermal degradation, thermal stability.

1. Introduction The thermal behavior of liquid-crystalline polyesters (LCPs) has been subject of some papers due to their unique technological properties, such as low melting viscosity, self reinforcing mechanical properties, and high flame resistance[1-4], which is verified by extensive char formation.

The improvement of the mechanical properties of commodity and engineering thermoplastics upon addition of a small amount of liquid-crystalline polyester has been reported, even in the absence of miscibility in most of those systems[5,6].

Increased compatibility in polymer blends has been reported when a flexible liquid-crystalline polyester was employed and interfacial adhesion was observed[7,8]. In recent years, block copolymers containing liquid-crystalline segments and isotropic segments have been suggested to represent the frontier in this kind of polymer blends in terms of modulation of stiffness and size domains[9,10].

Unlike LCPs, under thermal degradation conditions, most of thermoplastics and lightly crosslinked elastomers go through heat induced transformations that include hydrogen stripping, side group elimination, depolymerization to monomers and main chain rupture, followed by complete volatilization with no residue left. The influence of the presence of LCPs on the thermal degradative processes of thermoplastics is not well-known, especially in blends with isotropic-anisotropic segmented LCPs. Thermogravimetric analysis (TGA) can supply information on the nature of the degradative reactions, if competitive or independent, and occasionally should be an useful tool to cast some light on the interactions present in a sample.

The thermal degradative process of aromatic polyesters containing a flexible segment has been already described.

The degradation proceeds by free radical mechanism after the primary breakage of the ester linkage of the terephthalic acid (Figure 1). This mechanism involves a cyclic transition state with the hydrogen at the β position, followed by a chain rupture process in which vinyl ester groups and carboxylic

280

acids are formed[11-13] (Figure 1). A second stage was reported to be the thermal scission of the mesogenic group with phenylenic residues formation. The final degradation stage corresponds to the decomposition and volatilization of the rigid segments.

In this paper, the thermal stability of a series of copolymers comprised by polystyrene and liquid-crystalline oligoester segments is studied. The copolymers present liquid crystalline properties and a nematic mesophase is observed by polarizing optical microscopy and thermal analysis and they are distinguished by the following general chemical structure (Figure 2).

The thermal degradative process of polystyrene is also previously described in many texbooks and a radical mechanism with random chain scission and chain terminal unzipping or depropagation is involved in the main process[14].

McNeill et al.[14] demonstrated by a series of degradation studies of polystyrene, that between 200 and 300 °C, the molecular weight of polystyrene is decreased without formation of volatile products, and above 300 °C, polystyrene degraded to products which can be separated into two main fractions: (1) volatile products comprising mainly monomer and small amounts of toluene and α-methylstyrene, and (2) volatile products under vacuum and degradation temperature, the so-called cold ring fraction (CRF), comprising dimer, trimer and other short chain fragments. The secondary PS macroradical was confirmed to be the main source of volatile and CRF products. After 25% volatilization, intramolecular radical transfer becomes the main reaction and the proportion of CRF products is higher.

In our study, block copolymers and polymer blends comprised by the corresponding homopolymers, polystyrene and liquid-crystalline oligoester, were submitted to thermal analysis in order to understand the whole thermal degradative process.

Polímeros, 27(4), 280-284, 2017

Thermal degradation of polymer systems having liquid crystalline oligoester segment

2. Materials and Methods

2.3 Characterization

2.1 Synthesis of block copolymers

The chemical composition of the block copolymers (samples 2-5) was determined by 1H-NMR using a Varian Gemini 200 spectrometer.

Thermotropic block copolymers, poly(methyl-1,4dioxiphenylene-4,4’-dicarbonyl-1,10-dibenzoyl-decane)b-polystyrene, were prepared by a three step synthetic process[15-17]. In a first step a dihydroxy-terminated oligoester (sample 1, Table 1) was prepared from a polycondensation reaction of 1,10-dibenzoyl-decane-4,4’-dicarboxylic acid in the presence of a molar excess of methyl-hydroquinone. In a second step, the dihydroxy-terminated oligoester formed was further esterified with azo-ciano-pentanoyl chloride in order to produce a free radical macroinitiator. The block copolymers (samples 2-5, Table 1) were obtained in a third step by radical polymerization of styrene carried out in the presence of the macroinitiator.

2.2 Preparation of blends Polymer blends (samples 7-9, Table 1) were prepared by dissolution of oligoester (sample 1) and polystyrene Enichem PS1380 in chloroform (sample 6), followed by coprecipitation into methanol. Precipitated mixtures were vacuum dried at room temperature for 12 hours.

The block copolymers (samples 2-5) were hydrolyzed to yield the correspondent polystyrene blocks (Table 2) by dispersing the copolymer (0.5 g) in methanol (70 ml) together with sodium hydroxide (7.0 g). The complete hydrolysis (3-4h) of the mixture was monitored by IR spectroscopy. The polystyrene residue was filtered, washed with methanol and dried.

The molar masses of oligoester, polystyrene and block copolymers were determined by gel permeation chromatography (GPC) in THF, using an LC10 Perkin-Elmer liquid chromatograph equipped with a Jasco 830-RI refractive index detector and 103 Å and 104 Å Waters ultrastyragel columns. Monodisperse polystyrene standards were used for calibration.

2.4 Thermal analysis Thermogravimetry was carried out with a Mettler TG50 thermoanalyser, under dynamic nitrogen atmosphere of 50 ml/min flow and heating rate of 10°C/min. Sample size was 10 to 15 mg.

DSC analysis were performed with a Mettler DSC-30 calorimeter under nitrogen atmosphere and heating and cooling cycles of 10 °C/min. Sample size was 10 to 20 mg and temperature range was -20°C to 250°C.

3. Results and Discussions

Figure 1. Initiation mechanism of terephthalic polyester decomposition containing a flexible segment[11-13].

Average molar masses of the polystyrene segments varied with the copolymer composition. Higher molar mass polystyrene observed in the block copolymers containing lower concentrations of oligoester, were due to the synthetical process adopted[16].

Figure 2. Block copolymer polystyrene-b-polyester.

Table 1. Homopolymers, copolymers and blends prepared in this work.

Sample

Type

1 2 3 4 5 6 7 8 9

homopolymer copolymer copolymer copolymer copolymer homopolymer blend blend blend

Oligoester

Polystyrene

(%) 100 63 39 29 16 0 70 50 30

(%) 0 37 61 71 84 100a 30a 50a 70a

Comercial polystyrene PS1380 used in the blends.

a

Polímeros, 27(4), 280-284, 2017

The relative composition of the two polymeric segments in the block copolymers was determined by 1H-NMR spectrometry (Table 2). The integration correspondent to the signal due to the benzene ortho-hydrogens, at 6.5 ppm, was used for polystyrene content determination, and the integration of the methylene oxide hydrogens signal, at 4.4 ppm, for the estimation of the oligoester.

DSC analysis of these materials confirmed the existence of a stronger intersegmental interaction in the copolymers than in the polymer blends (Figure 3). A decrease of the Tg of the polystyrene phase in both polymeric systems, block copolymers and polymer blends, was observed and it is attributed to a partial miscibility of the oligoester in the polystyrene phase.

In the block copolymers, the crystalline phase also changed.

Broadening and shift of the melting (Tm) and isotropization (Ti) peaks were observed. The lowering of the isotropization temperature of oligoesters is atributted to the solvation of crystallites by polystyrene segments. The extension of the Ti variation has been shown to be dependent on the size of both polymeric segments[16]. Blends of oligoester with the commercial polystyrene did not present any shift on the melting or isotropization temperature.

281

Matroniani, R., & Wang, S. H. Table 2. Average molar masses and compositions of block copolymers, homopolymers, and recovered polystyrene blocks.

Sample 1 2 3 4 5 6a

Oligoester (%) 100 63 39 29 16 -

Mw

Block copolymer Mw/Mn

4500 10000 21000 54500 110500 -

1.61 1.72 2.22 1.52 1.48 -

Mw

Polystyrene block Mw/Mn

8000 17500 28500 69500 228000

1.39 1.73 1.71 1.63 1.47

Comercial polystyrene PS1380 used in the blends.

a

Figure 3. Transition temperature, Tg and Ti, in block copolymers and polymer blends as a function of composition. Ti in polymer blend (■), Ti in block copolymer (◻), Tg in polymer blend (●) e Tg in block copolymer (○).

These results clearly indicate an extensive miscibility in the polystyrene amorphous phase, in the block copolymer or polymer blend. However, a separated oligoester phase was still present as isotropization temperature was observed.

Although constant in the polymer blends, and indicating essentially no significant intermolecular interactions, Ti in copolymers steadily decreased as the polystyrene segment length decreased from Mw=28500 g/mol to Mw= 8000 g/mol, suggesting again segmental interaction within this oligoester phase.

The oligoester, the block copolymers and the polymer blends under the conditions of the thermogravimetric analysis presented two main weight losses (Figure 4). The polystyrene presented only one sharp weight loss, primarily owing to the depropagation mechanism, as described in literature (Figure 4). The thermal decomposition of flexible thermotropic polyester based on a mesogenic triad unit with three aromatic rings has been described to occur in two fundamental steps in the path to a carbonaceous residue[17]. According to the authors[17], the first step is initiated by fragmentation of the flexible segment and vaporization of dienes, however immediately before the second step these fragments are no longer detected by mass spectroscopy and elimination of CO2 is observed after reaching 400 °C, with a second maximum rate around 460 °C, as observed for sample 1, oligoester (Figure 4).

Table 3 presents the results of TGA as a function of copolymer composition.

282

Figure 4. TGA of the homopolymers, sample 1 (oligoester 1) and polystyrene Enichem PS1380, and blend of oligoester 1 and polystyrene.

At the onset of the degradation, Td1, at very low weight loss, the degradation mechanism is still governed by intermolecular radical transfer. This temperature reflects the intermolecular interactions and is dependent on the molar mass of the sample.

Polímeros, 27(4), 280-284, 2017

Thermal degradation of polymer systems having liquid crystalline oligoester segment Table 3. TGA of block copolymers prepared from oligoester 1.

Sample 1 2 3 4 5

%

Td1

wt-loss-1

Td2

wt-loss-2

residue

oligoester 100 63 39 29 16

°C 353 352 336 338 361

% 62 81 68 77 100

°C 426 402 412 419 -

% 23 23 32 23 0

% 15 0 0 0 0

Td1 temperature at the first onset; Td2 temperature at the second onset; wt-loss-1 weight loss in the first stage; and wt-loss-2 weight loss in the second stage.

Table 4. TGA of blends of polystyrene 6 and oligoester 1.

Sample 6 7 8 9

%

Td1

wt-loss-1

Td2

wt-loss-2

residue

oligoester 0 70 50 30

°C 372 335 349 361

% 100 71 86 84

°C 417 415 413

% 0 16 12 13

% 0 13 2 3

Td1 temperature of first curve inflection; Td2 temperature of the second deflection; wt-loss-1 weight loss in the first stage; and wt-loss-2 weight loss in the second stage.

A lowering in the thermal stability in both copolymers and polymer blends was observed, as the first weight loss derivative curve exposed a shoulder around 360°C (exemplified by sample 8, indicated in Figure 4) and a total weight loss in the first stage over 70%. The residue after heating to 600°C in copolymers and blends were also lower than that observed for the oligoester.

Such behavior suggests that a degradative process of the aromatic segments of the oligoester was induced by the radicals formed during the fast polystyrene decomposition.

The relative higher Td1 observed for sample 2 compared to other samples, copolymers and blends, is unexpected because of the higher miscibility observed in both crystalline and amorphous phases (Figure 3). For this sample, a thermal degradation due to intermolecular radical transfer seems to be limited, maybe because of its comparatively lower molar mass (Table 2).

Samples 3 and 4 with comparable polystyrene segment lengths and miscibility behaviors (Table 2 and Figure 3), were also similar in thermal induced decomposition and both degradation steps presented comparable onset temperatures.

Accordingly to its low oligoester content, sample 5 showed the expected higher thermal stability compared to 3 and 4.

The dependence of the degradation temperature of polystyrene on its molar mass is well known, higher stability is expected for the higher molar mass samples. However, in the polymer blends, the molar masses of the components were kept constant and the decomposition behavior should be only ascribed to the composition of each blend.

Table 4 shows that the onset of degradation was dependent on the oligoester concentration when molar masses were kept constant. The increase in oligoester concentration decreased the thermal stability of those blends.

A comparison between copolymer 4 and blend 9 shows that the same concentration of oligoester is present, however, the thermal stability of copolymer 4 was much lower. The lower thermal stability observed for copolymer Polímeros, 27(4), 280-284, 2017

should be attributed to the existence of stronger segmental interaction and lower molar mass, reinforcing a mutual degradative process.

4. Conclusions Compared to the homopolymers, LC-oligoester and polystyrene, the block copolymers and the polymer blends having liquid crystalline segments presented a lower thermal stability. The lower thermal stability observed is attributed to the existence of an interaction between the two types of polymer chain, polystyrene and oligoester. It is suggested that in these polymer systems, under thermal induced degradative conditions, transport of radical species between phases favored the degradation propagation. The even higher instability of the copolymers compared to blends is believed to be a result of lower molar mass and stronger segmental interaction, as it was observed by DSC analysis, causing an intensified mutual polymer degradation process.

However, the judicious incorporation of LCP to commodity polymers may be a potential procedure to achieve standard thermoplastics with incremented thermal resistance, as char formation is common in those systems.

5. Acknowledgements The authors thank Brazilian CAPES and CNPq for financial support.

6. References 1. Jin, X., & Chung, T. S. (1999). Thermal decomposition behavior of main-chain thermotropic liquid crystalline polymers, vectra A-950, B-950, and Xydar SRT-900. Journal of Applied Polymer Science, 73(11), 2195-2207. http://dx.doi. org/10.1002/(SICI)1097-4628(19990912)73:11<2195::AIDAPP17>3.0.CO;2-3.

2. Mandal, P. K., Siddhanta, S. K., & Chakraborty, D. (2011).

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Matroniani, R., & Wang, S. H. liquid crystalline polymer blends. Journal of Applied Polymer Science, 124(6), 5279-5285.

3. Wei, P., Wang, L., Wang, X. H., Chen, Y. W., Wang, Y. P., & Wang, Y. M. (2014). Nonisothermal and isothermal oxidative degradation behavior of thermotropic liquid crystal polyesters containing kinked bisphenol AF and bisphenol A units.

High Performance Polymers, 26(8), 935-945. http://dx.doi. org/10.1177/0954008314535645.

4. Chi, Z. G., Yao, X. D., Zhang, Y., & Xu, J. R. (2005). Thermal decomposition kinetics of thermotropic liquid crystalline polyesterimides. Journal of Applied Polymer Science, 98(6), 2467-2472. http://dx.doi.org/10.1002/app.22447.

5. Mithal, A. K., Tayebi, A., & Lin, C. H. (1991). In situ composite fibers: Blends of liquid crystalline polymer and poly (ethylene terephthalate). Polymer Engineering and Science, 31(21), 1533-1538. http://dx.doi.org/10.1002/pen.760312105.

6. Mehta, S., & Deopura, D. L. (1993). Fibers from blends of PET and thermotropic liquid-crystalline polymer. Polymer Engineering and Science, 33(14), 931-936. http://dx.doi. org/10.1002/pen.760331410.

7. Shin, B. Y., & Chung, I. J. (1990). Polymer blend containing a thermotropic polyester with long flexible spacer in the main chain. Polymer Engineering and Science, 30(1), 22-29. http:// dx.doi.org/10.1002/pen.760300105.

8. Shin, B. Y., Jang, S. H., Chung, I. J., & Kim, B. S. (1992).

Mechanical-properties and morphology of polymer blends of poly(ethylene-terephthalate) and semiflexible thermotropic liquid-crystalline polyesters. Polymer Engineering and Science, 32(1), 73-79. http://dx.doi.org/10.1002/pen.760320112.

9. Heitz, T., Rohrbach, P., & Höcker, H. (1989). Rigid rods with flexible side-chains - a route to molecular reinforcement.

Macromolecular Chemistry and Physics, 190(12), 3295-3316. http://dx.doi.org/10.1002/macp.1989.021901226.

10. Kobayashi, T., Sato, M., Takeno, N., & Mukaida, K. (1992).

Synthesis and liquid crystallinity of thermotropic polycarbonatepolystyrene graft-copolymers. European Polymer Journal, 28(9), 1105-1110. http://dx.doi.org/10.1016/0014-3057(92)90062-7.

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11. Lum, R. M. (1979). Thermal-decomposition of poly(butylene terephthalate). Journal of Polymer Science. Part A, Polymer Chemistry, 17(1), 203-213. http://dx.doi.org/10.1002/ pol.1979.170170120.

12. Foti, S., Giuffrida, M., Maravigna, P., & Montaudo, G. (1984).

Direct mass-spectrometry of polymers. 11. Primary thermal fragmentation processes in aromatic aliphatic polyesters.

Journal of Polymer Science. Part A, Polymer Chemistry, 22(6), 1217-1229. http://dx.doi.org/10.1002/pol.1984.170220602.

13. Montaudo, G., Puglisi, C., Scamporrino, E., & Vitalini, D. (1986). Thermal degradation of aromatic-aliphatic polyethers.

1. Direct pyrolysis-mass spectrometry. Macromolecules, 19(3), 870-882. http://dx.doi.org/10.1021/ma00157a067.

14. McNeill, I. C., Zulfiqar, M., & Kousar, T. (1990). A detailed investigation of the products of the thermal-degradation of polystyrene. Polymer Degradation & Stability, 28(2), 131-151. http://dx.doi.org/10.1016/0141-3910(90)90002-O. 15. Wang, S. H., Coutinho, F. M. B., Galli, G., & Chielllini, E. (1996). Copolímeros em bloco termotrópicos poliestireno – Poli(Metil-1,4-Dioxifenileno-4,4-Dicarbonil-1,10-DibenzoilOxi-Decano): 1. Síntese e caracterização. Polímeros: Ciência e Tecnologia, 6(4), 38-44.

16. Wang, S. H., Coutinho, F. M. B., Galli, G., & Chiellini, E. (1995). Synthesis and characterization of polystyrene-polyester liquid-crystalline block-copolymers. Polymer Bulletin, 34(5-6), 531-537. http://dx.doi.org/10.1007/BF00423348.

17. Ellis, G., Marco, C., del Pino, J., & Goméz, M. A. J. (1998).

Thermal stability and degradation mechanism for two mainchain liquid crystal polyesters - A TG-MS study. Journal of Thermal Analysis and Calorimetry, 52(3), 683-695. http:// dx.doi.org/10.1023/A:1010185817984.

Received: Dec. 23, 2015 Revised: Apr. 18, 2016 Accepted: June 29, 2016

Polímeros, 27(4), 280-284, 2017

http://dx.doi.org/10.1590/0104-1428.2375

New technologies from the bioworld: selection of biopolymer-producing microalgae Roberta Guimarães Martins1, Igor Severo Gonçalves2, Michele Greque de Morais2 and Jorge Alberto Vieira Costa1* Laboratório de Engenharia Bioquímica, Escola de Química e Alimentos, Universidade Federal do Rio Grande – FURG, Rio Grande, RS, Brazil 2 Laboratório de Microbiologia e Bioquímica, Escola de Química e Alimentos, Universidade Federal do Rio Grande – FURG, Rio Grande, RS, Brazil 1

*jorgealbertovc@terra.com.br

Abstract Microalgae are studied because of their biotechnological potential. The growth of microalgae aims at obtaining natural compounds. Due to the large amount of accumulated polymer waste, one of the solutions is the use of biodegradable polymers. The objective of this work was to select biopolymer-producing microalgae and to study the cell growth phase in which maximum production occurs. Microalgae Cyanobium sp., Nostoc ellipsosporum, Spirulina sp. LEB 18 and Synechococcus nidulans were studied. The growth was carried out in closed 2 L photobioreactors kept in a chamber thermostated at 30 °C with an illuminance of 41.6 μmolphotons.m-2.s-1 and a 12 h light/dark photoperiod. The biopolymers were extracted at times of 5, 10, 15, 20 and 25 d. The microalgae that had the highest yields were Nostoc ellipsosporum and Spirulina sp. LEB 18 with crude biopolymer efficiency of 19.27 and 20.62% in 10 and 15 d, respectively, at the maximum cell growth phase.

Keywords: cyanobacteria, biopolymer, polyhydroxyalkanoate, productivity.

1. Introduction Cyanobacteria were the first phototrophic organisms capable of producing oxygen. They are responsible for the conversion of Earth’s atmosphere from anoxic to oxic[1].

For the production of biomass with specific characteristics, manipulation of the culture conditions is a key factor[2].

Cyanobacteria are used for various purposes, e.g., for food supplements for humans[3] and animals[4]. Some cultures are used in wastewater treatment[5], in fixing carbon dioxide and in biocompound synthesis[6,7]. The biomass of Spirulina has been investigated for its hypocholesterolemic potential[8], as a source of biofuels[9] and for biopolymer production[10-12]. Several genera and species of cyanobacteria, such as Dunaliella tertiolecta[11], Aulosira fertilissima[12], Nostoc muscorum[13], Spirulina subsalsa[14], Synechocystis sp.[15], Spirulina platensis[16] and Synechococcus sp.[17], are used for the production of biopolymers.

Bacteria and cyanobacteria have the capacity to produce polyhydroxyalkanoates (PHAs)[13,18], which are biodegradable polyesters with potential use as polymeric materials[19].

Biodegradable polymers are alternative replacements for petrochemical polymers[20].

Reducing the consumption of plastic materials is difficult because of their versatile properties. However, it is possible to replace the petrochemical polymers with alternative materials that have similar polymer properties but show rapid degradation after disposal[20].

PHAs may positively change the scenario of global climate impact by reducing the amount of non-biodegradable polymers used[20]. Mixed cyanobacterial and bacterial cultures to produce PHAs are emerging due to the potential

Polímeros, 27(4), 285-289, 2017

residuary use for growth and low installation cost towards a profitable production of polyhydroxyalkanoates. The growth of microalgae does not require large amounts of land and can occupy areas unsuitable for agriculture, thus not competing with food production, due to the possibility of using photobioreactors that maximize biomass production[21,22].

The objective of this work was to select biopolymer‑producing microalgae and to study the phase of cell growth in which maximum production occurs.

2. Materials and Methods 2.1 Microorganisms and culture medium The microalgae used were Cyanobium sp., Nostoc ellipsosporum, Spirulina sp. LEB 18 and Synechococcus nidulans. The microalgal strain Nostoc ellipsosporum (B1453-79) was provided by the University of Göttingen (Germany). The cyanobacteria Cyanobium sp.[23], Spirulina sp. LEB 18[24] and Synechococcus nidulans[7] belong to the Collection of Strains of the Laboratory of Biochemical Engineering of the Federal University of Rio Grande (FURG). Spirulina sp. LEB 18 was isolated from Mangueira Lagoon (33°30’12” S, 53°08’58” W) located in Santa Vitoria do Palmar/RS (Brazil). The cyanobacterium Synechococcus nidulans was isolated from a stabilization pond of the President Medici Thermoelectric Power Plant, located in Candiota/RS (Brazil) (24º36’13”S, 52º32’43”W).

Inocula of Cyanobium sp. and Nostoc ellipsosporum microalgae were maintained in BG-11 culture medium[25], and Spirulina sp. LEB 18 and Synechococcus nidulans

285

O O O O O O O O O O O O O O O O

Martins, R. G., Gonçalves, I. S., Morais, M. G., & Costa, J. A. V. microalgae were maintained in Zarrouk culture medium[26].

All inoculations were adapted to their respective culture media for 30 d before the start of the experiments.

2.2 Culture conditions The cultivations were performed in closed 2 L photobioreactors with a working volume of 1.5 L and continuous agitation by the injection of sterile air to avoid the precipitation of the biomass. For Nostoc ellipsosporum, Spirulina sp. LEB 18 and Synechococcus nidulans, the initial concentration was 0.15 g.L-1, but for Cyanobium sp., the initial concentration was 0.2 g.L-1. The triplicate cultures were kept in a thermostated chamber at 30 °C for 5, 10, 15, 20 and 25 d, for a total of 15 experiments for each microalgae. The illuminance used was 41.6 μmolphotons.m-2.s-1 with a 12 h light/dark photoperiod maintained by 40 W fluorescent lamps.

2.3 Analytical determinations Daily samples were collected aseptically for the monitoring of the cell concentration and pH. Cell concentration was determined by optical density at 670 nm in a spectrophotometer (Quimis Q798DRM, Brazil) with a calibration curve relating the optical density to the dry weight of the microalgal biomass[27].

The pH determination was performed in digital pH meter (Quimis Q400H, Brazil) following AOAC methodology[28].

2.4 Determination of the crude biopolymer yield The crude biopolymer yield (YCB) was calculated according to Equation 1, where Ccbt is the concentration of crude biopolymers (g.L-1) at time t (d), Ccb5 is the concentration of crude biopolymers (g.L-1) at time 5 d, t is the time (d), and t5 is the time at 5 d.

YCB = (Ccbt - Ccb5 ) / (t - t 5 ) (1)

2.5 Extraction of crude biopolymers After 5, 10, 15, 20 and 25 d of experiment, the cultures were centrifuged at 7500 rpm for 20 min at room temperature (Hitachi, Japan) to separate the wet biomass from the biopolymer of the culture medium. Later, for every 1 g of dry biomass, 100 mL of distilled water and 25 mL of sodium hypochlorite (10-12% active chlorine (w/v)) were added to the wet biomass, and the solution was kept under stirring for 10 min. The resulting suspension was centrifuged (7500 rpm for 20 min at room temperature). Then, the supernatant was discarded, and the precipitate was washed with 100 mL of distilled water. The sample was centrifuged again, and the supernatant was discarded. This process was repeated adding 50 mL of acetone. The final precipitate (crude biopolymers) was dried at 35 °C for 48 h. The efficiency (η) of crude biopolymers in relation to microalgal biomass (%) was calculated using Equation 2, where mcb is the final mass of crude biopolymer obtained from the microalgal biomass (g), and mma is microalgal biomass (g). η = (mcb *100) / m ma (2) 286

2.6 Statistical analysis The results were processed by analysis of variance (ANOVA) and Tukey’s test to compare the means of the parameters analyzed with a 95% confidence level.

3. Results and Discussions The growth curves of cyanobacteria Cyanobium sp., Nostoc ellipsosporum, Spirulina sp. LEB 18 and Synechococcus nidulans (Figure 1) showed different behaviors in spite of each species having its own specific growth characteristics and different culture media. In preliminary tests, it was observed that when the microalga Cyanobium sp. was grown at low biomass concentrations (0.15 g.L-1), it showed photoinhibition in its growth; therefore, the assays were carried out with an initial biomass concentration of 0.2 g.L-1, thereby preventing cell death and providing the lag phase of growth.

Spirulina sp. LEB 18 (Figure 1c) showed early stationary growth phase after 20 d of culture. For Cyanobium sp., Nostoc ellipsosporum and Synechococcus nidulans, the stationary phase of growth was not observed by the end of the 25 d of culture. To verify the growth phases of the microalgae Cyanobium sp., N. ellipsosporum and S. nidulans, it would be necessary to grow the cultures for a longer period. For large-scale production, such a long culture period is impractical for the production of biopolymers. Sharma and Mallick[29] cultivated Nostoc muscorum microalgae in BG-11 medium with a phosphorus deficiency and addition of exogenous carbon sources and found an increase in the production of PHB. Yields of up to 8.6% (PHB) were found when the extraction of the polymer was performed in the early stationary phase of growth of the microalgae (21 d of culture), whereas in log phase, the yield was 6.1%. Samantaray and Mallick[12] cultivated the microalga Aulosira fertilissima during 14 d and observed an accumulation of 6.4% of PHB at the end of logarithmic growth phase.

The microalga Nostoc ellipsosporum presented a different behavior in its cell growth compared to the other microalgae under study. During the first 8 d of culture, it showed cell growth, then ceased and remained constant until the 17th d, after which it presented new cell growth.

This growth pattern may have occurred because when the microalgae are under a particular nutrient limitation, they use a substrate from its own cell as a nutrient, enabling continued growth. If there is a lack of carbon, the microorganism can consume the biopolymer itself. In this case, it is believed that the biopolymer may have been consumed, because after the 10th d of cultivation, the yield of biopolymers was reduced (Table 1). Another nutrient that may have had an influence was nitrogen, whose release in the culture medium from amino acids of phycobiliproteins and chlorophyll can possibly allow cell maintenance to occur[30,31].

The cyanobacterium Nostoc ellipsosporum presented a cell concentration less than the others but had higher efficiency (Table 1) and crude biopolymer yield (Table 2).

Among the microalgae under study, Nostoc ellipsosporum and Spirulina sp. LEB 18 stood out. These microalgae showed the higher efficiency of crude biopolymers (PHB) and did not differ significantly (p<0.05) each other from 15 d.

However, Nostoc ellipsosporum reached a crude biopolymer Polímeros, 27(4), 285-289, 2017

New technologies from the bioworld: selection of biopolymer-producing microalgae

Figure 1. Growth curves of microalgae Cyanobium sp. (a) Nostoc ellipsosporum (b), Spirulina sp. LEB 18 (c) and Synechococcus nidulans (d) with 5 (■), 10 (●), 15 (▲), 20 (♦) and 25 (+) d of culture.

Table 1. Crude biopolymer efficiency (%, w/w*) for microalgae at different culture times.

Time (d) 5 10 15 20 25

Cyanobium sp.

3.68±0.23aAB 3.17±0.26abA 2.75±0.40bA 2.91±0.15abA 3.12±0.30abA

Microalgae N. ellipsosporum Spirulina sp. LEB 18 9.04±3.24aC 5.82±2.02aB 19.27±1.18bB 10.23±0.93aC 17.79±1.32bB 20.62±3.17bB abB 13.41±3.80 11.83±1.67aB 10.69±2.84aB 11.86±2.43aB

S. nidulans 1.18±0.23aA 8.83±0.06bC 1.00±0.33aA 10.21±1.95bcB 11.01±1.49cB

For the same letters, the averages do not differ significantly (p<0.05) by Tukey test. Lowercase letters compare the results in columns. Uppercase letters compare the results in the rows. *Values correspond to averages of results obtained in triplicate with their respective standard deviations.

Table 2. Crude biopolymer yield (Ycb, gcb.L-1.d-1) for microalgae at different culture times.

Time (d) 5 10 15 20 25

Cyanobium sp. <0.01 <0.01 <0.01 <0.01

Polímeros, 27(4), 285-289, 2017

Microalgae Nostoc ellipsosporum Spirulina sp. LEB 18 2.05 0.88 0.87 1.48 0.29 0.40 0.08 0.30

Synechococcus nidulans 1.53 <0.01 0.60 0.49

287

Martins, R. G., Gonçalves, I. S., Morais, M. G., & Costa, J. A. V. efficiency of 19.27% ​​in 10 d and Spirulina sp. LEB 18 reached 20.62% in 15 d of culture. The crude biopolymer efficiency of Nostoc ellipsosporum was 2.05 g.L-1.d-1 at 10 d, where as that of Spirulina sp. LEB 18 was 1.48 g.L-1.d-1 at 15 d (Table 2). Panda et al.[15] found that the cyanobacterium Synechocystis sp. PCC 6803 accumulated biopolymer PHB in its cells. It has been found that when cultured in BG-11 medium under phosphorus and/or nitrogen deficiency with the addition of exogenous carbon sources, this microalgae showed a higher yield (4.5%) of PHB in the early stationary growth phase (at 21 d cultivation), while in the logarithmic phase, the yield was 2.9%.

The microalga Cyanobium sp. did not achieve significant results (p>0.05) for the production of crude biopolymers.

The cyanobacterium Synechococcus nidulans showed the highest PHB efficiency (11.01±1.49%) at a greater time of growth (25 d) in relation to the microalgae Spirulinasp.

LEB 18 and Nostoc ellipsosporum. Therefore, its use is less interesting compared to Nostoc ellipsosporum and Spirulina sp. LEB 18. Lower yields (3%) of PHB were found

by Sankhla et al.[32] in the stationary phase of growth when studying the production of PHB by Brevibacillus invocatus MTCC 9039.

The lowest yields obtained in culture times greater than 10 d (Nostoc ellipsosporum) and 15 d (Spirulina sp.

LEB 18) may be due to the depletion of nutrients from the medium, especially carbon, which leads to consumption of the biopolymers for cell growth and maintenance. The results showed the effect of culture time on the production of biopolymers. This difference in yield is associated with the fact that the production of the polymer depends on the availability of the source of carbon and energy, which vary as a function of the culture time. Bhati and Mallick[13] studied the microalga Nostoc muscorum for the production of PHB-HV with yields of 16.6% in 10 d of incubation.

For the same microalga, yields of different biopolymers were observed at different times using different carbon sources. When BG-11 medium was used with the addition of propionate, the highest yield was 12.6% in 21 d and 16.6% in 10 d with the addition of valerate. The highest yields were in the late exponential phase of growth. Mallick[33] studied the production of PHB-HV in Nostoc muscorum using BG-11 medium with the addition of propionate yielding 28.2% of biopolymer in 14 d of culture (late exponential growth phase).

Several microalgae, especially cyanobacteria, are able to accumulate intracellular biopolymers, especially poly-3-hydroxybutyrate and poly (3-hydroxybutyrateco-3-hydroxyvalerate) belonging to the group of polyhydroxyalkanoates. By modifying the culture conditions, particularly the nutrients, one can divert the metabolic pathways, causing the microorganism to synthesize larger amounts of biopolymers.

Studies are being carried out with photosynthetic mixtures of bacteria and algae that accumulate PHA in conditions with different concentrations of nutrients, and these studies have achieved PHB yields of 20%. The use of mixed photosynthetic culture (bacteria and microalgae) has emerged as an alternative system for the production of 288

PHA, potentially minimizing feed costs through the use of solar energy[34].

The defatted biomass of microalgae Dunaliella tertiolecta was used for the production of biopolymers in different salt concentrations, obtaining a yield of 82%[11]. High yields of biopolymers can be achieved using microalgae. It is possible to conclude that many microalgae are able to intracellularly accumulate PHB granules. However, different behaviors are observed due to the use of different microalgal sources and concentrations of nutrients and growth conditions.

4. Conclusions This study showed that in order to produce biopolymers from microalgal cultures, the microalgae Spirulina sp. LEB 18 and Nostoc ellipsosporum would be the best candidates.

Both microalgae had higher concentrations of biopolymers at short growth times (Spirulina sp. LEB 18, 20.62% in 15 d; Nostoc ellipsosporum, 19.27% in 10 d). Combining the growth of microalgae and biopolymer production is a strategy with the potential to significantly reduce environmental pollution problems, through both the use of industrial waste as a source of nutrients for the culture medium and the replacement of petrochemical origin polymers by biopolymers degradable and compostable when disposed of in the environment.

5. Acknowledgements The authors would like to thank CNPq (National Council of Technological and Scientific Development), CGTEE (Company of Thermal Generation of Electric Power) and MCTI (Ministry of Science, Technology and Inovation) for their financial support of this study.

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32. Sankhla, I. S., Bhati, R., Singh, A. K., & Mallick, N. (2010).

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33. Mallick, N., Gupta, S., Panda, B., & Sen, R. (2007). Process optimization for poly(3-hydroxybutyrate-co-3-hydroxyvalerate) co-polymer production by Nostoc muscorum. Biochemical Engineering Journal, 37(2), 125-130. http://dx.doi.org/10.1016/j. bej.2007.04.002.

34. Fradinho, J. C., Domingos, J. M. B., Carvalho, G., Oehmen, A., & Reis, M. A. M. (2013). Polyhydroxyalkanoates production by a mixed photosynthetic consortium of bacteria and algae.

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Received: Nov. 09, 2015 Revised: Mar. 29, 2016 Accepted: May 17, 2016 289

http://dx.doi.org/10.1590/0104-1428.2403

O O O O O O O O O O O O O O O O

Microalgae biopeptides applied in nanofibers for the development of active packaging Carolina Ferrer Gonçalves1, Daiane Angelica Schmatz1, Lívia da Silva Uebel1, Suelen Goettems Kuntzler1, Jorge Alberto Vieira Costa2, Karine Rigon Zimmer1 and Michele Greque de Morais1* Laboratory of Microbiology and Biochemistry, College of Chemistry and Food Engineering, Universidade Federal do Rio Grande – FURG, Rio Grande, RS, Brazil 2 Laboratory of Biochemical Engineering, College of Chemistry and Food Engineering, Universidade Federal do Rio Grande – FURG, Rio Grande, RS, Brazil

1

*michele.morais@pq.cnpq.br

Abstract This study was conducted to develop PCL nanofibers with the incorporation of microalgae biopeptides and to evaluate the stability of chicken meat cuts during storage. PCL and PCL/biopeptides nanofibers were formed by electrospinning method, and the diameters obtained were 404 and 438 nm, respectively. The tensile strength, elongation, melting temperature and thermal stability of biopeptide-added PCL nanofibers were 0.245 MPa, 64%, 56.8 °C and 318 °C, respectively. PCL/biopeptide nanofibers showed a reducing power of 0.182, inhibition of 22.6% and 12.4% for DPPH and ABTS radicals, respectively. Chicken meat cuts covered by the PCL/biopeptide nanofibers showed 0.98 mgMDA∙kg-1 and 25.8 mgN∙100g-1 for TBARS and N-BVT analysis, respectively. Thus, the PCL/biopeptide nanofibers provided greater stability to the product and control of oxidative processes ensuring the product quality maintenance during the 12 d of storage.

Keywords: antioxidants, electrospinning, poly-ɛ-caprolactone.

1. Introduction The conservation of fresh meat products is an important factor in ensuring food safety to the final consumer. Due to the lipid and protein content, these products are targets of lipid oxidation that result in nutritional changes by the degradation of fat-soluble vitamins and essential fatty acids[1,2]. New technologies have been employed to improve the quality and extend the shelf life of food products. The use of active packaging is one such technology and consists in the action of antioxidants and antimicrobial compounds that interact with food[3,4]. There are several mechanisms of action of these agents that include absorption of carbon dioxide, oxygen, ethylene and odors. Furthermore, compounds such as antimicrobials and antioxidants that retard the degradation processes in food have been used in packaging[5-7].

In the food industry, many synthetic antioxidants such as butyl hydroxy toluene (BHT) and butyl hydroxy anisole (BHA) are used to slow the peroxidation processes. However, the use of these compounds must be controlled due to the carcinogenic effects on human health[8]. Thus, the search for natural antioxidants is a safer alternative for use in food.

Microalgae are capable of synthesizing many bioactive compounds. These include lipids, carotenoids and phycobiliproteins. The biopeptides of microalgal source have applications as dietary supplements, health promoters and more recently they have been suggested for inclusion in active packaging[9]. Spirulina is a cyanobacterium that has GRAS (Generally Recognized As Safe) certification, with high protein content and is a source of biopeptides with antioxidant activity[10]. The biopeptides are protein

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fragments that contain 3 to 20 amino acid residues that are inactive within the protein molecule and can be released by hydrolysis[11]. The application studies of biopeptides of microalgal source are restricted in the packaging area for food preservation[12].

The nanofibers can be applied to food preservation.

Packaging formed by nanofibers have advantages for allowing an increase of the contact area of t​​ he product with bioactive compounds. The incorporation of biopeptides in the matrix of nanofibers has the aim of inferring improvements in performance and in the physical and active properties of carriers. Nanofiber packages are alternatives to increase food shelf life to the end consumer[13].

Nanofibers can also be developed from biodegradable polymers. This reduces the environmental problems caused by the disposal of packaging developed from polymers of petrochemical origin[14-16]. In this sense, the objective of this study was to develop nanofibers with the incorporation of antioxidant biopeptides of microalgal source for the conservation of chicken meat.

2. Materials and Methods 2.1 Obtaining biopeptides by enzymatic hydrolysis of the Spirulina sp. LEB 18 biomass Spirulina sp. LEB 18 biomass was obtained from the pilot plant of the Laboratory of Biochemical Engineering, located in the city of Santa Vitória do Palmar, Rio Grande do Sul[17]. The biomass was concentrated in a hydraulic

Polímeros, 27(4), 290-297, 2017

Microalgae biopeptides applied in nanofibers for the development of active packaging press, dried at 50 °C, ground in a ball mill (QUIMIS Q298), sieved (ABNT/Tyler 60) and kept at -18 °C. Protein hydrolysis was carried out in 100 mL reactors using 3% of the microalgae biomass solubilized in sodium carbonate bicarbonate buffer pH 9.5 and 3 U∙mL-1 of Protemax 580 L enzyme, courtesy of Prozyn (São Paulo).

The process was conducted at 60 °C under agitation of 180 rpm for 240 min. The final reaction was heat inactivated at 85 °C bath for 10 min. The degree of hydrolysis (DH) was determined by the protein content before and after the process, according to the method described by Hoyle and Merrit[18]. The protein hydrolysates were filtered through qualitative membranes of 0.45 μm, 0.22 μm and 0.1 μm and Amicon 10K vertical column. After this step the samples were lyophilized.

2.2 Development of nanofibers

The melting temperature was determined from the peak shown in the DSC melting curve[20].

The thermal stability of nanofibers, and residual solvent was carried out in thermogravimetric analyzer (Shimadzu DTG-60, Japan) according to ASTM D3850-12 protocol[21].

Analysis were conducted from 25 to 500 °C under an inert nitrogen atmosphere with a flow rate of 30 mL∙min-1 and constant heating rate of 10 °C∙min-1 using 3 mg of sample.

The tensile strength and elongation at break of the nanofibers were measured by a texturometer (Stable Micro Systems Model TA.XT plus, England). Samples were prepared with dimensions of 10 x 70 mm and thickness measured in a micrometer (Starrett 444MXRL-75, Brazil). Assays were performed at speeds of 2 mm∙s-1 and initial distance between grips of 50 mm.

The tensile strength and elongation at break were calculated according to Equations 1 and 2.

F Ts = m (1) A where: Ts = tensile strength (MPa); Fm = maximum force at the time of rupture of the nanofibers (N); A = cross‑sectional area (m2). d ε =100 r (2) di

Polycaprolactone (PCL) polymer obtained from Sigma Aldrich (density of 1.145 g∙mL-1 and molecular weight of 80,000 g∙mol-1) was used in preparing polymer solutions for the development of nanofibers. The solution contained 12% (w/v) PCL, 1.4% (w/v) NaCl, 3% (w/v) biopeptides using chloroform:methanol (1:3, v/v) as solvent for the solubilization of the compounds in the polymeric solution. The control solution was prepared under the same conditions, containing only PCL and NaCl. The solutions were homogenized in a where: ɛ = elongation at break (%); dI = initial separation magnetic stirrer for 12 h (25 °C). distance (mm); dR is the difference between the separation The PCL and biopeptide-added PCL solutions from distance at the time of rupture and the initial distance. microalgal source were placed in syringe with capillary The antioxidant activity was determined for the of 0.70 mm diameter and injected across infusion pump Spirulina sp. LEB 18 biomass, biopeptides as well as the (KD Scientific, KDS 100, USA). The potential difference PCL and biopeptide-added PCL nanofibers were filtered between capillary and collector caused evaporation of the through Amicon 10K column. The methods evaluated solvent and the nanofibers were deposited on the collector. the reducing power[22] and sequestration capacity of the The distance between the capillary and collector was free radical DPPH (2,2-diphenyl-1-picryl-hidrazol)[23], and 120 mm, electric potential of 25 kV, and solution feed rate of -1 2000 μL∙h . The process environment condition was 25 °C ABTS (2,2’-azino-bis(3-ethylbenzothiazoline-6-sulfonic and relative humidity 44%. The nanofibers were formed acid))[24].The measurements were expressed as inhibition using a solution volume of 2 mL. After the process, the percentage (Equation 3). nanofibers were collected and stored in a desiccator under  ABSblank − ABS sample  controlled humidity (20% R.H). = % Inhibition   ×100 (3)

2.3 Evaluation of developed nanofibers Analysis were performed on samples of PCL (control) and biopeptide-added PCL nanofibers. The nanofibers were analyzed in a scanning electron microscope (SEM) (JEOL JSM-6610 LV, Japan). The diameters were determined using 30 readings of nanofibers. The samples were fixed in a metallic support and coated with gold using diode sputtering (Denton Vacuum CAR001-0038, USA) according to ASTM E986-04[19].

The viscosity of the polymeric solutions was determined by rheometer (Brookfield Programmable DV-III Ultra Rheometer, USA). This analysis consists of the direct measurement of the viscosity of PCL and biopeptide-added PCL solutions.

The melting temperatures and enthalpies were determined by analysis of differential scanning calorimetry (DSC) (Shimadzu DSC-60, Japan). A sample of 3 mg of nanofibers was placed under nitrogen atmosphere and flow of 50 mL∙min-1. The analysis were conducted at range between 25 °C and 180 °C, at heating rate of 10 °C∙min-1.

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

ABSblank



The biomass solution and biopeptides were prepared at a concentration of 10 mg∙mL-1 for analysis of the reducing power methods and sequestering of ABTS free radical.

In DPPH free radical sequestration analysis, solutions were prepared at a concentration of 5 mg∙mL-1. For the nanofibers, 50 mg of the nanofiber samples were solubilized with the addition of 5 mL methanol and 2 mL of chloroform for the rupture of the structure and extraction of biopeptides.

The true concentration of biopeptides in the matrix of nanofibers analyzed for antioxidant activity corresponded to 0.21 mg∙mL-1. The solutions were homogenized by vortex for 1 min.

2.4 Application of bioactive nanofibers and evaluation of the stability of chicken meat cuts during storage In the analysis of stability, cuts of the same chicken breast sample of approximately 90 g each were done. A sample was coated with the nanofiber matrix containing biopeptides and another sample was left without coverage for 12 d, the 291

Gonçalves, C. F., Schmatz, D. A., Uebel, L. S., Kuntzler, S. G., Costa, J. A. V., Zimmer, K. R., & Morais, M. G. validity period for commercial poultry. Samples were then cut and homogenized for analysis.

The stability of chicken breasts stored under refrigeration (± 6 °C) with the nanofibers containing biopeptides and without nanofibers was evaluated by the test of reactive species to the 2-thiobarbituric acid (TBARS) with modifications[25].

Chicken breast samples with a mass of 50 g were cut out and homogenized with 100 mL of 7.5% trichloroacetic acid (TCA) for 20 min in a mixer, vacuum filtered and the volume completed to 100 mL in a volumetric flask. A 5 mL aliquot of the filtrate was mixed with 5 mL of 0.02 M thiobarbituric acid (TBA) in test tube covered with an aluminum foil and placed on a water bath for 30 min at 80 °C. The analysis of the blank containing 5 mL of 7.5% TCA and 5 mL of TBA was carried out parallel to the assays. Soon after, the reading was carried out by spectrophotometer at 538 nm.

The TEP (tetraethoxypropane) standard curve was used for the quantification of TBARS. The evaluations were performed in two stages: at time zero and after 12 d of storage. Nanofiber matrices were removed from chicken samples for analysis.

The determination of total volatile bases (N-BVT) it was made in chicken cuts kept under refrigeration (± 6 °C)[26].

A sample of approximately 50 g was blended with 100 mL of 7.5% TCA for 20 min on a mixer, vacuum filtered and the volume completed to 100 mL in a volumetric flask.

An aliquot of 10 mL of the extracts was transferred to micro Kjeldahl distillation tube, 3 drops of phenolphthalein were added and it was subjected to distillation. The distillate was collected in 5 mL of boric acid (50 g∙L-1) with 4 drops of bromocresol green and methyl red indicator (30:20).

The titration was performed with a 0.02 N hydrochloric acid solution. The calculation was performed according to Equation 4, and expressed in mgN∙100g-1 sample.

(

)

VHClsample − VHClblank × N HCl ×14.01×100 (4) N − BVT = Psample

where in: VHCl= volume (mL) used in the titration; NHCl = normality of HCl; Psample = mass of sample (g).

2.5 Statistical analysis Analysis were performed in triplicate and the results were evaluated by analysis of variance (ANOVA) one-way at the 95% level of confidence.

3. Results and Discussion The apparent viscosities of the PCL and biopeptide‑added PCL solutions were 221.2 ± 5.1 mPa∙s-1 and 243.1 ± 12.0 mPa∙s-1, respectively, and made the formation of cylindrical nanofibers possible without forming droplets. The small increase in viscosity can be due to the presence of methanol, which according to product specifications shows higher viscosity (600 mPa∙s-1) than chloroform (580 mPa∙s-1). Besides that, this increase on viscosity is related with polarity of chloroform present in greater proportion on solution, making it difficult to solubilize the biopeptide which is a polar molecule. These solution characteristics are important for the electrospinning process because they prevent the formation of drops and allow the continuous jet of polymer solution to form nanofibers with greater uniformity[27].

PCL 10% (w/v) solutions used by Ranjbar-Mohammadi and Bahrami[28] showed apparent viscosity ranging from 700 mPa∙s-1. The authors developed nanofibers by electrospinning with a distance of 150 mm between the capillary and collector, electric potential of 15 kV and flow rate of 2000 μL∙h-1, presenting an average diameter of 156 nm[28]. The values in the present study are in agreement, the difference in viscosity can be associated with the addition of methanol for solubilization of the biopeptides in the PCL solution. The addition of methanol helped to increase the polarity solution, however, the added methanol fraction was less than chloroform and therefore did not provide complete homogenization of the solution, that can also have increased viscosity.

PCL and biopeptide-added PCL nanofibers produced by electrospinning were observed by SEM (Figure 1) to verify the form and the average diameter. Cylindrical nanofibers were obtained in a nanometer scale with diameter

Figure 1. Nanofibers with 12% PCL, 1.4% NaCl (a) and 12% PCL, 1.4% NaCl, methanol and 3% Spirulina sp. LEB 18 biopeptides (b) under magnification of 3500x.

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Microalgae biopeptides applied in nanofibers for the development of active packaging uniform, which can be confirmed by the standard deviations.

The diameters of the nanofibers were 404 ± 72 nm and 438 ± 24 nm for PCL and biopeptide-added PCL respectively.

The diameter of nanofibers containing the biopeptides did not differ significantly at 95% confidence in relation to the PCL nanofibers.

This study showed similar results to that developed by Goes et al.[29]. The authors obtained nanofibers from 15% PCL solution with addition of 2.5% clay and resulted in an average diameter of 340 nm[29]. Furthermore, the solutions were prepared by mixing chloroform and methanol as in this study. Thus, the conductivity characteristics were improved due to the increased polarity of the solution by the presence of methanol and salt. Wu et al.[7], obtained PCL fibers with the addition of 20% (w/w) polyaniline without formation of droplets with a diameter of 150 nm. The solutions were also prepared with a mixture of chloroform and methanol.

easily at high temperatures, different from what happens with PCL that is a synthetic polymer with higher thermal stability.

Therefore, the biopeptides addition in PCL nanofibers caused the decrease on degradation temperature. Furthermore, there was no degradation peak in the range of 60 °C, however, it is found that no residue of chloroform and methanol solvents are in the nanofibers, being completely evaporated after the electrospinning process.

The changes in the temperatures of degradation of nanofibers containing biopeptides may be a consequence crystallinity change of the PCL caused by addition natural compounds that are degraded at lower temperatures. Ciardeli et al.[32] also observed similar behavior with PCL/polysaccharides blends reported reduction in temperature in the pyrolysis compared to pure PCL films. Patrício et al.[33] obtained thermal stability temperatures values of ​​ up to 300 °C for polymer blends of PCL and PLA.

According to DSC analysis data, only one transition peak was presented for the PCL and biopeptide-added PCL samples (Figure 2). The DSC analysis resulted in a first-order endothermic event that may result from processes such as breaking bonds, decomposition and volatilization. Table 1 shows the melting point values (T ​​ melting) and enthalpy for PCL and biopeptide-added PCL. The enthalpy is closely related to the amount of energy absorbed by the samples for the change of state to occur.

Campos et al.[30] evaluated the thermal properties of extruded PCL films and reported that the Tmelting was 56.36 °C being in accordance with the values of ​​ this study.

In addition, they report that the glass transition temperature of PCL is in the range of –60 °C. From this information, it is possible to set the applicability range of PCL as food packaging for chilled products or at room temperature, since the temperatures do not exceed the limits of the change of state of nanometer material.

In a study by Wang et al.[31] it was found that changes in Tmelting in PCL nanofibers alter the crystallinity of the polymer and consequently the biodegradation process.

Moreover, they observed that the solvent is evaporated lasting through the electrospinning process and even after deposition of the nanofibers on the collector and the residual solvent continues to evaporate. Thus, the use of nanofibers as food packaging becomes secure. In a study of PCL blends associated with polysaccharides, Ciardeli et al.[32] obtained a reduction in Tmelting values of ​​ the samples when compared to the PCL film due to the interactions of the compounds incorporated into the polymer.

In the derivative curves (Figure 3), the biopeptide‑added PCL nanofibers showed (Table 1) ​​ initial and final thermal degradation temperature smaller compared to PCL nanofibers.

This might have occurred, because of the biopeptides sensibility which, being natural compound, degrade more

Figure 2. DSC curves for nanofiber PCL and biopeptide-added PCL.

Figure 3. TGA curves for nanofiber PCL and biopeptide-added PCL.

Table 1. Thermal properties of DSC and TGA of PCL and biopeptide-added PCL nanofibers.

Sample

Tmelting (ºC)

Enthalpy (J.g-1)

Tid (ºC)

Tm (ºC)

Tfd (ºC)

PCL PCL + biopeptides

58.6 56.8

34.9 28.6

333.0 262.3

403.4 318.0

434.0 339.6

Tmelting: melting temperature of the nanofibers; Tid: initial degradation temperature of nanofibers; Tm: average degradation temperature of nanofibers; Tfd: Final degradation temperature of nanofibers.

Polímeros, 27(4), 290-297, 2017

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Gonçalves, C. F., Schmatz, D. A., Uebel, L. S., Kuntzler, S. G., Costa, J. A. V., Zimmer, K. R., & Morais, M. G. The changes in thermal parameters evaluated by DSC and TGA analysis are not harmful when obtaining nanofibers aimed for application in packaging. Thus, one can define that the processing of nanofibers and their applicability should be performed below 260 °C. Obtaining biopeptide-added nanofibers by electrospinning is suitable because it does not use high temperatures in the process. Thus, biopeptides are not denatured and the maintenance of the activity for application as packaging occurs. The information about the thermal parameters are important to define the possible applications, as well as recycling and disposal. Moreover, the stability characteristics of nanofibers have potential for food packaging and the degradability that biopolymers are beneficial to have reduced environmental problems[34].

Mechanical properties of PCL and biopeptide-added PCL nanofibers were evaluated. The thickness of the films were 0.50 and 0.16 mm for PCL and biopeptide-added PCL nanofibers, respectively. This difference is due to the nanofibers were randomly produced in flat collectors without alignment, and have not been disposed homogeneously in the collector occurred the emergence of areas with the greatest amount of nanofibers, which probably caused the values for the tensile strength presented inferior results to those found in other studies.

The tensile strength values ​​are dependent on the maximum force applied to the material for the rupture to occur and cross-sectional area of ​​the matrix of nanofibers.

The tensile strength and elongation differences are directly linked to the thickness of the matrix of nanofibers. PCL and biopeptide-added PCL nanofibers showed tensile strengths of 0.137 MPa and 0.245 MPa, and elongation of 85% and 64%, respectively. Johnson et al.[35] obtained tensile strength of 1.29 MPa and elongation of 102% to nanofibers formed from 12% PCL solution. This value was higher than that obtained in this study because the authors used acetone under heating (50 °C) for the preparation of the polymer solution.

The change of the interaction between solvent and polymer may form a distinct organization in the formed polymer chains which directly influences the properties of nanofibers.

In a study by Ghasemi-Mobarakeh et al.[36] was obtained tensile strength of 3 MPa to PCL nanofibers and when was added gelatin in PCL nanofibers was observed that there was a reduction in tensile strength,resulting in less than 1 MPa. This occurred due to the gelatin is a natural polymer that presents inferior mechanical properties compared to PCL nanofibers resulting in less mechanical strength than pure PCL nanofiber. This same event was observed in the present study when biopeptide-added in nanofibers.

For the DPPH method, there was no significant difference of the inhibition percentage between unhydrolyzed biomass and biopeptide-added nanofibers (Table 2). The ABTS

sequestration methods and reducing power showed significant differences for the samples analyzed, showing increased antioxidant activity of the compounds after hydrolysis (Table 2). The increased activity of the peptides in relation to biomass was expected, due to be produced via intracellular by the biomass may have masked their activity.

The reduction of antioxidant activity after the electrospinning process was observed in all methods.

The reduction may be associated with losses during the process of production das nanofibras and also on obtaining of the extracts for quantification of antioxidant activity.

Likewise, the nanofiber matrix is composed ​​ of only 3% (w/v) biopeptides corresponding to 0.21 mg∙mL-1 of the compound in the extract while in the analysis of the pure compound, 5 mg∙mL-1 and 10 mg∙mL-1 of biopeptides were used. Still, biopeptides contained in PCL nanofibers showed antioxidant activity for all three methods studies.

Sheih et al.[37] in a study to obtain biopeptides from seaweed residues obtained compounds of a molecular mass of 1.3 kDa and tested the activity against DPPH and ABTS methods. The authors obtained 50% sequestration of these radicals at low concentrations of approximately 10 μg∙mL-1 when compared with synthetic antioxidants.

Cian et al.[38] studied the bioactivity of purified peptides obtained by hydrolysis of the algae Porphyracolumbina and showed values of 50% inhibition of DPPH and ABTS radicals at concentrations of approximately 3 mg∙mL-1. The antioxidant activity of biopeptides obtained from Spirulina sp. LEB 18 were significant, since at low concentrations it was possible to obtain non-purified biopeptides with antioxidant activity against the tested methods.

Figure 4 shows the chicken meat samples with application of nanofiber matrix with biopeptides and the control. The control sample without nanofiber coating showed significant difference (p <0.05) in the content of malondialdehyde and total volatile bases compared to chicken meat samples stored with nanofibers. These figures show that the chicken meat without coating showed higher production of compounds derived from lipid oxidation (Table 3).

The application of nanofibers with antioxidant biopeptides in chicken meat was efficient during storage and reduced degradation of the sample. Nanofibers have greater surface area contact compared to their polymeric counterparts in macroscopic scale. Thus, the bioactive compounds showed higher reactivity with the degradation products aiding in product quality maintenance[15].

The use of nanocomposites in the food packaging industry is promising since it greatly enhances the shelf life of products such as meats, cheeses, fruits and cereals[39].

In this study it is observed that the application of nanofibers

Table 2. Antioxidant activity of biomass, biopeptides and PCL and biopeptide-added PCL nanofibers.

Sample Biomass Biopeptides PCL PCL + biopeptides

DPPH (%Inhibition) 28.3 ± 3.8ª,b 30.6 ± 1.1a ̶ ** 22.6 ± 2.9b

ABTS (%Inhibition) 26.5 ± 1.7b 58.3 ± 0.9a ̶ ** 12.4 ± 0.6c

Poder Redutor (U.A., λ= 700 nm) 0.415 ± 0.015b 0.677 ± 0.007a 0.006 ± 0.001d 0.182 ± 0.024c

Different letters in the same column represent statistically different results (p <0.05); ̶ **: absorbance values ​​were used as a blank for the calculation of % inhibition of DPPH and ABTS radicals.

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Microalgae biopeptides applied in nanofibers for the development of active packaging

Figure 4. Control chicken cuts (1) and with nanofibers (2) containing biopeptides after 0 (a) and 12 d (b) of storage.

Table 3. TBARS analysis and N-BVT for samples containing nanofibers with biopeptides and the control after 12 d of storage.

Sample Control (0 d) Control (12 d) PCL+ biopeptides (12 d)

TBARS(mgMDA∙kg-1) N-BVT (mgN∙100g-1) 0.3 ± 0.0c 25.1 ± 0.7b 2.6 ± 0.0a 33.8 ± 0.0a 1.0 ± 0.0b 25.8 ± 0.3b

Different letters in the same column represent statistically different results (p <0.05).

with biopeptides maintained the conservation of the product during the storage period evaluated while the sample without nanofiber coating showed values above those considered normal for maintaining the organoleptic characteristics of chicken meat.

Counsell and Horning[40]reported that at TBARS values above 2 mg malondialdehyde per kg sample, rancid odors are detected by untrained judges. Bazargani-Gilani et al.[41] studied the stability of chicken breast coated with chitosan enriched with antioxidant plant extracts during storage under refrigeration. The application of natural antioxidants controlled oxidative processes during storage presenting potential for conservation of samples.

When coated with nanofibers containing biopeptides, chicken meat presented control in the process of degradation and values below those set as a limit by law. The quantification of volatile bases is generally regulated for fish. There is Polímeros, 27(4), 290-297, 2017

no current legislation to determine the N-BVT content of chicken meat cuts. Some studies show that levels above 30 mgN∙100g-1 give the product sensory changes making it to be considered as not suitable for human consumption[42].

4. Conclusions The nanofibers containing biopeptides of microalgal origin presented a diameter of 438 nm and due to the large surface area of contact can be applied in food preservation. In the assays of stability of chicken cuts with the application of nanofibers containing biopeptides, there was control in the lipid oxidation process with values of 1.0 ± 0.0 mgMDA∙kg-1 and 25.8 ± 0.3 mgN∙100g-1 compared with the control sample which showed 2.6 ± 0.0 mgMDA∙kg-1 and 33.8 ± 0.0 mgN∙100g-1. Thus, PCL nanofiber matrices containing biopeptides with antioxidant activity are potential alternatives for use as primary active packaging with the aim of conservation of food products.

5. Acknowledgements The authors would like to thank CAPES - Coordenação de Aperfeiçoamento de Pessoal de Nível Superior, CNPq – Conselho Nacional de Pesquisa Científica e Tecnológica, MCTI – Ministério da Ciência, Tecnologia e Inovação and CEME-SUL – Centro de Microscopia Eletrônica do Sul for their financial support for this study.

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O O O O O O O O O O O O O O O O

Modification of thermal and rheological behavior of asphalt binder by the addition of an ethylene-methyl acrylate-glycidyl methacrylate terpolymer and polyphosphoric acid Gerson da Silva Pereira1 and Ana Rita Morales1* Department of Materials Science and Bioprocess, School of Chemical Engineering, Universidade Estadual de Campinas – UNICAMP, Campinas, SP, Brazil

1

*morales@feq.unicamp.br

Abstract This study evaluated the modification effects of adding ethylene-methyl acrylate-glycidyl methacrylate terpolymer (EMGMA) in the presence of polyphosphoric acid (PPA) to an asphalt binder graded as 50/70 (0.1mm) in the Brazilian penetration grade specification (AC 50/70). The EMGMA terpolymer has been presented as a new alternative to modify asphalt binders properties, as scientific literature is scarce on its usage in this context and also on the role of PPA when used in combination with reactive polymers. The characteristics of the modified binder were analyzed by standard and rheological testing, including Multiple Stress Creep Recovery test (MSCR) and Fourier Transform Infrared Spectroscopy (FTIR) analysis. The MSCR test showed that the modified binder presented lower values of non-recoverable compliances (Jnr) and a higher percent recovery, when compared to the conventional binder. This behavior indicates that addition of EMGMA and PPA in asphalt binders could enhance the resistance to rutting of asphalt mixtures. The statistical evaluation showed that EMGMA had greater influence on the studied properties of Jnr (0.1kPa), MSCR recovery, softening point and elastic recovery at 25°C and that the PPA had also significant influence on these properties. FTIR analysis showed that chemical reactions occurred between the asphalt binder and EMGMA, forming a three-dimensional polymeric network, which promotes improved characteristics.

Keywords: asphalt binder, modified asphalt binders, MSCR, non-recoverable compliance, ethylene-methyl acrylate-glycidyl methacrylate terpolymer.

1. Introduction Modified asphalt binders have been increasingly used in paving applications as an interesting alternative to deal with shortcoming of the petroleum asphalt cement (PAC).

Nowadays, the use of modified asphalt binders is an excellent alternative for more durable flexible pavements, improving some deficiencies and enhancing asphalt concrete properties such as resistance to rutting and cracking and reducing the damage due to fatigue and thermal susceptibility[1].

Several polymeric modifiers may confer beneficial characteristics to asphalt binders, among them can be cited styrene-butadiene block copolymer (SBR), styrene-butadiene-styrene (SBS), styrene-isoprene-styrene (SIS), styrene-ethylene-butadiene-styrene (SEBS), acrylonitrile-butadiene-styrene (ABS), ethylene vinyl acetate (EVA)[2]. SBS copolymer is considered the most appropriate material for asphalt modification, despite the addition of it has economic limitations and may have serious technical limitations, as the low resistance to heat and oxidation[1].

Reactive terpolymers have been also used as modifiers of asphalts binders. Their reactivity is due to the presence of functional groups that are able to interact with the carboxyl groups of the asphaltenes. The asphaltenes fraction consists of large micelles with average molecular weight estimated between 800-3500 g/mol, and each micelle has a large number of available carboxyl groups, as shown in Figure 1.

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The terpolymers reacts with the asphaltenes and develops a chemical bond network between micelles improving the rheological properties of the asphalt binder. Also, there is the epoxy opening ring that can react with the hydroxyl groups and causes the formation of an ether bond. The same reaction may happen through an amino group, and once the epoxy ring is open, crosslinking between the polymer chains may occur, not necessarily involving asphaltene molecules[3]. Examples of reactive polymers are thermoplastic elastomers functionalized with maleic anhydride and ethylene copolymers containing epoxy rings[3]. Various studies show the improvement of the characteristics of PAC by adding ethylene-butyl acrylate-glycidyl methacrylate terpolymer (EBGMA)[3-6]. Studies with EBGMA terpolymer show important aspects to be considered in these systems, for instance, the polymer concentration, the curing time and the curing temperature that can promote a substantial increase in viscosity, which may induce the system to gel[3]. According to standard tests of penetration, softening point, rotational viscosity and aging, the modified asphalt binder with EBGMA became more consistent and with less thermal susceptibility compared to the unmodified asphalt binder[4]. It was also reported that the addition of 1% EBGMA terpolymer improved the rheological properties of the binder, providing greater resistance to rutting[5].

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Modification of thermal and rheological behavior of asphalt binder by the addition of an ethylene-methyl acrylate-glycidyl methacrylate terpolymer and polyphosphoric acid Some studies also related the use of the polyphosphoric acid (PPA) in very low amounts to catalyze the reaction between the terpolymer and the asphalt molecular groups, as this is a slow reaction that requires up to 24 hours to complete the cure[3,7]. The addition of PPA in modified asphalts can also decrease the amount of polymer, reducing costs[8,9].

Related to the rheological properties, a system of PAC + EBGMA + PPA showed a very high deformation recovery and a very low non-recoverable compliance (Jnr) values, as well presented more resistance when subjected to a sudden increase in stress level[5]. In other study the addition of EBGMA terpolymer into the PAC modified its rheological properties, evidenced by an increase in the dynamic viscosity and a change of the Newtonian behavior at different temperatures[6].

Currently, the growing demand for modified asphalt binders has encouraged search for alternative new polymeric

materials that meet the performance and economic viability requirements. Within this context, a new ethylene-methyl acrylate–glycidyl methacrylate terpolymer (EMGMA), studied in this paper, was recently introduced to the market, this modifier being the main object of this work. The main difference between the new EMGMA and the most studied terpolymers is the replacement of the butyl acrylate monomer by the methyl acrylate monomer[11]. The chemical structure of EBGMA and EMGMA are presented in Figure 2a and b.

In this study, the addition of PPA was also evaluated since several researchers have looked into the use of this acid also as a modifier, either by itself or associated to other modifiers[12-14]. As described by many researchers, adding PPA in the PAC increases the content of asphaltenes, because it acts in this fraction as a deflocculant and increases the contact area with the maltenic phase[15,16].

Considering that scientific literature on the use of EMGMA terpolymer as asphalt binder modifier is scarce and the PPA function is not clear when used in combination with reactive polymers, this paper presents a laboratorial study of the thermal and rheological behavior for an asphalt binder modified with EMGMA in the presence of PPA in order to understand the polymer and PPA action and the components interactions.

2. Materials and Methods 2.1 Materials One PAC from Replan – Petrobras Refinery (Paulinia, São Paulo, Brazil) was used in this study. This unmodified material is graded as 50/70 in the Brazilian penetration grade (AC50/70), indicating a penetration of 50 -70 (dmm), according to ASTM D5-06.

Figure 1. 2-Dimensional model of an asphaltene molecule[10].

The EMGMA terpolymer was produced by Arkema (Lotader AX8900) and supplied by Quimigel Indústria Comércio Ltd., containing 23.13% of methyl acrylate and 8.35% of glycidyl methacrylate, having a glass transition temperature (Tg) of -2°C. It was supplied as extruded granules.

Figure 2. Molecular structure of (a) ethylene-butyl acrylate-glycidyl methacrylate terpolymer and (b) ethylene-methyl acrylate-glycidyl methacrylate terpolymer (adapted from Brito et al.[11]).

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Pereira, G. S., & Morales, A. R. Polyphosphoric Acid (PPA) is an inorganic oligomer obtained by the condensation of the monophosphoric acid or by the hydration of phosphorus pentoxide (P2O5).

The PPA is typically a mixture of orthophosphoric acid with pyrophosphoric acid and triphosphoric acid and does not contain free water. It is produced and marketed based on its relative level of H3PO4 (phosphoric acid) or P2O5 (phosphorus pentoxide) in its content[16]. In this study was used the PPA at the concentration of 84% of P2O5, proportional to 116% of H3PO4, supplied by the company Prana Petroquímica Ltda.

2.2 Preparation of modified asphalt binder samples The unmodified PAC was preheated at a temperature of 175°C using an electric hot plate. Next, the EMGMA polymer was added to the PAC at the rate of 10g/min. An IKA stirrer (model RW20) with a propeller type R 1342 with four blades, was used to promote mixing at a speed of 1750 rpm.

The stirring was done for 30 min, until the complete dispersion of the polymer, then, the PPA was added, remaining under agitation for additional 90 min.

2.3 Experimental design It was used a factorial experimental design of 22, as shown in Table 1, according to the methodology described

by Neto et al.[17].

The statistical analysis of the data was performed using the software Statistica 10.0, with a confidence level of 95%. The factorial analysis was used to assess the effects and interactions of the variables in the traditional binder properties of softening point and elastic recovery at 25°C, and also in the non-recoverable compliance (Jnr) and in the percent of recovery, obtained at Multiple Stress Creep and Recovery (MSCR) test.

shear stiffness modulus (G*) represents high rigidity, while a small phase angle represents a greater elastic response (recoverable). An asphalt binder with these properties provides to an asphalt mixtures better resistance to rutting.

Thus, the G*/sinδ parameter was originally adopted using the Superpave[23] methodology to express the resistance of asphalt binders to rutting. After the adoption of G*/sinδ in the Superpave specification, various studies showed that it was not appropriated to evaluate modified asphalt binders[24,25] because of its inability to assess the elastic contribution provided by the addition of polymers to the asphalt binders. Thus, in this study, the recovery percentage and the non-recoverable compliance values (Jnr) were determined, in the Multiple Stress Creep Recovery test (MSCR), following the procedure established by ASTM D7405-10[26]. The MSCR test applies an oscillatory regime to characterize permanent deformation of asphalt binders.

This test was performed on the DSR equipment using unaged asphalt binder samples. In the MSCR test, a constant stress of 0.1 kPa for 1 second is applied on an asphalt sample of 1 mm thickness and then the load is removed, allowing the material to relax for 9 seconds. This cycle is repeated 10 times, without time intervals between cycles.

After these 10 cycles, the same procedure is repeated using a constant stress of 3.2 kPa. The recovery percentage is obtained from the ratio of the recoverable deformation and the total deformation, as shown in Equation 1. ( EC − E0 ) − ( ER − E0 )  .100 (1) R ( σ, N ) = EC − E0

Where R(σ,N) is the percent recovery under stress σ (0.1 kPa or 3.2 kPa) and N is the cycle of the creep and recovery (1 ≤ N ≤ 10). Figure 3 shows the deformation values used for the calculations.

The properties evaluated by the conventional binder test were: penetration at 25°C (ASTM D5-06[18]), ring-and-ball softening point (ASTM D36-06[19]) and elastic recovery by ductilometer – (ASTM D6084-13[20]. The rotational viscosity was measured using a Brookfield Viscometer RVDV-I Prime model with spindle No. 20, according to ASTM D4402-06[21], at temperatures of 135°C, 155°C and 177° and rotational speeds of 20 rpm, 50 rpm and 100 rpm, respectively. The rheological tests were performed using a dynamic shear rheometer (DSR) from TA Instruments (DHR-1 model), with a parallel plate geometry of 25mm diameter and deformation amplitude of 12%, controlled automatically by the equipment software. The dissipation modulus, storage modulus and phase angle were also determined on the same DSR equipment, according ASTM D7175-08 standard[22].

The DSR is used to measure asphalt properties at high and intermediate temperatures. High values of complex

Figure 3. Deformations E0, EC and ER for one creep and recovery cycle of the MSCR test.

Table 1. Experimental design 22 for AC 50/70, EMGMA and PPA. Independent variables Levels

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% EMGMA Polymer 1.0 2.0 1.0 2.0

% PPA 0.05 0.30

Sample nomenclature (1% EMGMA + 0.05% PPA) (2% EMGMA + 0.05% PPA) (1% EMGMA + 0.30% PPA) (2% EMGMA + 0.30% PPA)

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Modification of thermal and rheological behavior of asphalt binder by the addition of an ethylene-methyl acrylate-glycidyl methacrylate terpolymer and polyphosphoric acid The non-recoverable compliance is the ratio of the non-recoverable deformation and the stress applied, and is typically expressed as the inverse of the stress unit (Pa-1 or kPa-1), according to Equation 2.

ER − E0 J NR ( σ, N ) = (2) σ

Where Jnr(σ,N) is the non-recoverable compliance relative to the stress σ at the creep and recovery cycle N, (1 ≤ N ≤ 10 and σ = 0.1 kPa or 3.2 kPa). With all the results of R(σ,N) and Jnr(σ,N) obtained in the 10 creep and recovery cycles, their average results were calculated by means of the simple arithmetic average. Studies have demonstrated that the Jnr value has good correlations with rutting measurements on asphalt mixtures[27,28].

With the purpose to investigate the chemical reaction that may occur between the asphalt binder and the EMGMA, analysis were performed by Fourier Transform Infrared Spectroscopy (FTIR) on a spectrometer Thermo Scientific Nicolet 670 model, measured in transmittance mode, in the range of 4000-675cm-1.

3. Results and Discussion Table 2 shows the results for unmodified and modified asphalt binder samples produced in accordance to the experimental design. The results indicated that binder with a higher PPA content had lower penetration values. Significant increase of the softening point was observed for all modified asphalt binder samples, increasing from 49°C to 76°C for the sample with 2% EMGMA + 0.30% PPA, which shows a lower temperature sensitivity of the asphalt binder. In the elastic recovery test performed at 25°C, it was noted that the modified binders have up to 6 times elastic recovery, when compared to the unmodified binder. The increase of the elastic recovery and the softening point values even with low EMGMA and PPA levels indicates that the terpolymer promotes the formation of a polymeric network, capable to provide an improvement in the elasticity of the PAC,

as described also by Polacco et al.[3] using the terpolymer EBGMA. Table 3 and Figure 4 present the results of rotational viscosity at high temperatures. It is noted that the addition of EMGMA increased the rotational viscosity of the asphalt binder. Adding higher amounts of PPA to the modified binders for the same EMGMA content, the rotational viscosity also increased, which may be associated to the crosslinking process. This behavior was also observed by other researchers when EBGMA was used[3,6].

As observed from the softening point and the elastic recovery results, a polymeric network in the asphalt matrix may be developed, which also promotes an increase of the rotational viscosity, hindering the molecules flow.

The rotational viscosity is important to measure the consistency properties of asphalt binder, which are used as a guide for processing and storing temperatures. Asphalt binders excessively viscous can be inappropriate to be used in asphalt mixtures.

In Brazil, a maximum rotational viscosity is specified for modified asphalt binders at each temperature, as shown in Table 3[29]. Therefore, as shown by the experimental data in Figure 4, only the modified asphalt binder with 2% EMGMA + 0.30% PPA exceeded the 3000 cP viscosity limit, indicating the possibility of operational problems.

The data obtained at the laboratory tests were analyzed using the Statistica software 10.0, at 95% confidence level and it was possible to evaluate the effects and interactions of the variables in the responses. This method is useful to analyze the results in a statistical base, pointing out the factors that are statistically significant. Figures 5 and 6 illustrates the effects and interactions of each factor (EMGMA % and PPA %) regarding the softening point and elastic recovery.

Figures 5 and 6 show that both, the EMGMA concentration as the PPA concentration, influence in the evaluated properties, with a statistically significance (p<0.05). The concentration of EMGMA was the most influential factor, because of its higher value in the softening point and elastic recovery

Table 2. Results of conventional tests of the AC 50/70 and the modified asphalt binder samples.

Sample AC 50/70 base binder (1% EMGMA + 0.05% PPA) (1% EMGMA + 0.30% PPA) (2% EMGMA + 0.05% PPA) (2% EMGMA + 0.30% PPA)

Penetration (dmm) 55+/-1 53+/-1 48+/-1 58+/-1 48+/-1

Softening point (°C) 49.0+/-0.8 55.0+/-0.6 62.0+/-0.3 66.0+/-0.5 77.0+/-0.8

Elastic recovery (%) 9.0+/- 0.3 71.0+/-0.8 77.0+/-0.8 83.0+/-1.0 92.0+/-0.8

Table 3. Rotational viscosity results and the maximum limits according to Brazilian specification.

Sample Spindle rotation speed Maximum limit[29] AC 50/70 base binder (1% EMGMA + 0.05% PPA) (1% EMGMA + 0.30% PPA) (2% EMGMA + 0.05% PPA) (2% EMGMA + 0.30% PPA)

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135°C 20 rpm 3000 345 625 937 1155 3840

Rotational viscosity (cP) 150°C 50 rpm 2000 158 304 388 450 1422

177°C 100 rpm 1000 63 110 145 140 343

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Pereira, G. S., & Morales, A. R. tests. Besides, there is a significant interaction between the factors in both responses, showing more pronounced in the softening point test. The evaluation of the response surface (Figure 5b and 6b) confirms that the concentration of EMGMA is the factor that most contributes to the increase of softening point and elastic recovery, but it is observed that the concentration of PPA also has strong influence on elastic recovery response. This suggests that PPA acts as a co-modifier of the PAC + EMGMA system, providing improved characteristics to the asphalt binder.

Figure 4. Rotational viscosities for the asphalt binder AC 50/70 and modified samples at temperatures of 135°C, 150°C and 177°C.

Conventional tests employed in the characterization of asphalt binder do not allow the determination of fundamental properties related to paved runway performance due to empiricism involved and practical complications related to the results interpretation[30]. To address this deficiency new tests based on rheological properties related to the performance of the asphalt mixture were established by the Superpave specification[25,26]. Thus, MSCR test was used to better assess the samples performance at performance temperatures of flexible pavements. The results of MSCR are shown in Figure 7 and it is noted that when EMGMA

Figure 5. Softening point (a) Pareto chart and (b) response surface.

Figure 6. Elastic recovery at 25°C (a) Pareto chart and (b) response surface.

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Modification of thermal and rheological behavior of asphalt binder by the addition of an ethylene-methyl acrylate-glycidyl methacrylate terpolymer and polyphosphoric acid

Figure 7. MSCR results for the applied stress of (a) 0.1 kPa and (b) 3.2 kPa.

content and PPA content are increased, the deformation decreases and the deformation recovery enhances.

The overall behavior of the samples, observed at MSCR test, was similar at both applied stresses levels (3.2 kPa and 0.1 kPa). However, as expected, higher deformation was obtained at the stress of 3.2 kPa. Only the sample with 1% EMGMA + 0.05%PPA showed a greater deformation compared to the AC50/70.

In Figure 8 are shown the 0.1 kPa and 3.2 kPa recovery results obtained by MSCR test and it can seen an increase in the deformation recovery because of the addition of EMGMA polymer. Samples with 1% EMGMA content exhibit significant recovery enhancement with increase in PPA addition. Samples with 2% EMGMA show the same behavior.

The effects and interactions of the MSCR recovery percentage at the stress of 0.1 kPa are illustrated in Figure 9.

It is noted that the EMGMA concentration has a high contribution in this property. Additionally, the lower PPA concentration significantly affects the response, but a higher percentage of PPA did not contribute to the increase of the MSCR recovery.

The non-recoverable compliance values (Jnr) at the stresses of 0.1 kPa and 3.2 kPa are shown in Figure 10.

Again, higher concentrations of EMGMA and PPA decrease the non-recoverable compliance values, except sample with 1% EMGMA and 0.05% of PPA. This behavior confirms the better performance of the modified binder, obtaining a EMGMA + PPA system that could contribute towards better rutting resistance of asphalt mixtures.

The statistical analysis of the effects and interaction of the variables on non-recoverable compliance (Jnr) response at the stress of 0.1 kPa stress is shown in Figure 11a and b.

Both variables (PPA% and EMGMA%) influence this response and the EMGMA concentration has a higher Polímeros, 27(4), 298-308, 2017

Figure 8. MSCR recovery percentage at stresses of 0.1 kPa and 3.2 kPa.

significance. The interaction between the two variables was more significant than the PPA concentration.

The MSCR test results obtained in this study using EMGMA are similar to those observed with EBGMA terpolymer, indicating that the mechanism of action in both cases are similar[5]. The improvement of the elastic characteristics with minimum percentage of EMGMA and PPA can only be explained by the formation of a three-dimensional polymer network formed by the bonding of epoxy groups EMGMA with asphaltene molecules[3].

The rheological properties of the samples were also studied in the linear viscoelastic region, defined as the strain interval where the complex modulus (G*) does 303

Pereira, G. S., & Morales, A. R.

Figure 9. MSCR recovery at the stress of 0.1 kPa (a) Pareto chart and (b) response surface.

behavior of the asphalt. The viscous dissipation, represented by loss modulus (G”) values indicate that, at relatively low temperatures (52°C), asphalts behave in a similar manner regarding the viscous component; however, with the increase of the temperature a gap is noted between unmodified asphalt binder and the modified samples. The viscoelastic behavior of the binder is also evaluated using the phase angle (δ).

At intermediate temperatures (normal conditions in asphalt mixtures), the asphalt binder has a viscoelastic behavior, and depending on the temperature and frequency, the peak response of asphalt binders may occur at any point for δ between 0 and 90°. This delay (δ) is a fundamental property to describe the viscoelastic behavior of asphalt binders.

Figure 10. Average of non-recoverable compliances (Jnr) at the stresses of 0.1 kPa and 3.2 kPa.

not vary with the applied shear stress, considering an accuracy of 5% for constant temperature and frequency[24].

Results of the complex modulus (G*) test were used to construct curves to better assess the modified samples.

Figure 12 shows the results, indicating that the modification process significantly extended the linear viscoelastic region for the samples (2% EMGMA + 0.05% PPA) and (2% EMGMA + 0.30% PPA). It is also observed a total overall linearity for the complex modulus to shear stress values between 90 and 150 Pa, that corresponds to the shear stress specified for the dynamic-mechanical tests for an asphalt binder without aging.

Figure 13 shows the viscoelastic behavior of modified asphalt binder, represented by the storage modulus (G’) and the loss modulus (G”) as function of the temperature.

As the EMGMA polymer and PPA percentage increase, the storage modulus (G’) also increases in all modified asphalt binder samples, indicating an improvement in the elastic 304

The higher the phase angle (δ) the higher the predominance of the viscous component. Figure 14 presents the δ values​​ of the samples studied. The unmodified asphalt binder had δ of 89° at 82°C showing a predominance of the viscous component. The sample (2% EMGMA + 0.30% PPA) had the lowest δ value, 55° at the same temperature, indicating a significant increase in the elastic properties of the modified asphalt.

Figure 15a shows the FTIR spectrum of the unmodified asphalt binder, EMGMA and the sample 2%EMGMA+0.3%PPA recorded in the range 2000-600 cm-1. The FTIR spectrum corresponding to EMGMA terpolymer shows a remarkable band at 1736cm-1, which is assigned to the carbonyl stretching (C=O) of glycidyl methacrylate (GMA)[31,32]. There are two bands of stretching vibration for ester compound: one is νas(C–O–C) at 1195 cm-1, the other is νs(C–O–C) at 1164 cm-1.

There is also a weak shoulder around 910cm-1 characteristic of the epoxy group[11,31,32].

Comparing the spectrum of unmodified asphalt binder and the modified one, the bands show few differences.

Therefore, it is not possible to identify any product of the reaction between the unmodified asphalt binder and EMGMA. Figure 15b presents a zoom that shows that the band characteristic of the epoxy group at 910cm-1 disappears in the modified asphalt binder. This indicates that the epoxy Polímeros, 27(4), 298-308, 2017

Modification of thermal and rheological behavior of asphalt binder by the addition of an ethylene-methyl acrylate-glycidyl methacrylate terpolymer and polyphosphoric acid

Figure 11. Non-recoverable compliance (Jnr) at the stress of 0.1 kPa (a) Pareto chart and (b) response surface.

Figure 12. Modulus complex G* as function of the shear stress for the samples tested at 64°C.

Figure 13. (a) Loss modulus (G”) and (b) storage modulus (G’) as function of the temperature.

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Figure 14. Phase angle (δ) as function of the temperature.

Figure 15. (a) FTIR spectrum of unmodified asphalt binder, pure EMGMA and the sample 2%EMGMA+0.3%PPA; (b) zoom for the epoxy groups.

rings of the EMGMA opened and reacted with the carboxyl groups in asphaltenes forming ester groups that are not differentiated from existing in the original material. As the epoxy groups EMGMA have a greater affinity reaction with the hydroxyl and carboxyl groups, and is more reactive with the carboxyl groups than with hydroxyl groups[11], the interaction between the terpolymer EMGMA and the PAC probably due the carboxylic groups that exist in abundance in the asphaltene fraction[3,10]. Moreover, many changes occur in the PAC because of the modification action of PPA that can be explained by the increase in the content of asphaltenes, as PPA acts as a deflocculant of this fraction, leading to an increase of the contact area of asphaltenes ​​ with maltenic phase[15,16]. Furthermore, at the beginning of the modification process, the dispersion of EMGMA and the reaction with 306

the carboxylic groups in the asphaltene fraction happens, but this reaction is slow[3,7]. With the addition of the PPA the deflocculation of asphaltene micelles occurs, which leverages the effect of EMGMA. Thus, a three-dimensional network which connects the asphaltene micelles is formed, as well as cross-reactivity with the polymer chain itself, contributing to the gain of the elastic properties of the asphalt binder.

4. Conclusion The addition of EMGMA terpolymer and the PPA into the asphalt binder AC 50/70 produced modified asphalt binders with improved elastic properties, as evidenced by the softening point tests and elastic recovery at 25°C. Polímeros, 27(4), 298-308, 2017

Modification of thermal and rheological behavior of asphalt binder by the addition of an ethylene-methyl acrylate-glycidyl methacrylate terpolymer and polyphosphoric acid The viscoelastic behavior showed lower phase angle values and higher storage modulus values.

The Multiple Stress Creep Recovery (MSCR) tests showed lower deformation and higher recovery stress values in the modified asphalt binders, besides, the results of non-recoverable compliances (Jnr) for applied stresses of 0.1 kPa and 3.2 kPa had a considerable decrease when EMGMA and PPA concentrations were increased. This fact is indicating that the modified asphalt binder may have a greater elastic response during the service conditions, contributing for a longer lifetime of the asphalt mixture in flexible pavements and a reduction in the susceptibility to rutting at longer loading-unloading times.

The statistical evaluation results give support to conclude that the EMGMA has greater influence on the non-recoverable compliance (Jnr) at 0.1 kPa, softening point and elastic recovery at 25°C. The PPA has also a significant influence on properties, meaning that it acts not only as a catalyst of the reaction between the asphalt components and the EMGMA terpolymer. There was also a significant interaction between the factors EMGMA% and PPA%.

FTIR spectrum showed that epoxy functional group of EMGMA disappear in the modified asphalt binder.

The improvement of the elastic characteristics with very low amount of EMGMA and PPA can only be explained by the development of a three-dimensional network formed by the reaction of epoxy groups EMGMA with the functional groups of the asphaltene molecules, while the PPA acts as a co-modifier.

The results presented agree to those observed with EBGMA terpolymer, indicating that the mechanisms of action in both cases are similar.

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ASTM D6084/D6084M – 13: Standard Test Method for Elastic Recovery of Bituminous Materials by Ductilometer. West Conshohocken: ASTM. 21. American Society for Testing and Materials – ASTM. (2006). ASTM D4402-06: Standard Test Method for Viscosity Determination of Asphalt at Elevated Temperatures Using a Rotational Viscometer. West Conshohocken: ASTM. 22. American Society for Testing and Materials – ASTM. (2008).

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Pereira, G. S., & Morales, A. R. E-C147). Washington: Transportation Research Board. Retrived in 2015, March 15, from http://onlinepubs.trb.org/onlinepubs/ circulars/ec147.pdf 24. Bouldin, M. G., Dongré, R., & D’Angelo, J. (2001). Proposed refinement of Superpave high-temperature specification parameter for performance-graded binders. Transportation Research Record, 1766, 40-47. http://dx.doi.org/10.3141/176606.

25. D’Angelo, J. (2015). Multi-stress creep and recovery test method a new specification. Austin: Association of Modified Asphalt Producers - AMAP. Retrived in 2015, March 15, from http://amap.ctcandassociates.com/wp/wp-content/uploads/ dangelo-MSCR-2-08E.pdf 26. American Society for Testing and Materials – ASTM. (2010).

ASTM D7405-10: Standard Test Method for Multiple Stress Creep and Recovery of Asphalt Binder Using a Dynamic Shear Rheometer. West Conshohocken: ASTM. 27. Hafeez, I., & Kamal, M. A. (2014). Creep compliance: a parameter to predict rut performance of asphalt binders and mixtures. Arabian Journal for Science and Engineering, 39(8), 5971-5978. http://dx.doi.org/10.1007/s13369-014-1216-2.

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29. Resolução ANP nº 32, de 21.9.2010. (2010, 22 de setembro).

Diário Oficial da República Federativa do Brasil, Brasília.

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32. Jun, L., Yuxia, Z., Yuzhen, Z. (2008). The research of GMAg-LDPE modified Qinhuangdao bitumen. Construction and Building Materials, 22, 1067-1073. http://dx.doi.org/10.1016/j. conbuildmat.2007.03.007.

Received: Dec. 23, 2015 Revised: May 31, 2016 Accepted: June 21, 2016

Polímeros, 27(4), 298-308, 2017

http://dx.doi.org/10.1590/0104-1428.00516

Dynamic mechanical analysis and crystalline analysis of hemp fiber reinforced cellulose filled epoxy composite Anand Palanivel1*, Anbumalar Veerabathiran2, Rajesh Duruvasalu1, Saranraj Iyyanar1 and Ramesh Velumayil1 1

Mechanical Engineering, Vel Tech Dr RR. & Dr. SR University, Avadi, Chennai, Tamilnadu, India 2 Mechanical Engineering, Velammal College of Engineering & Technology, Madurai, TN, India *p.anand@ymail.com

Abstract The Dynamic mechanical behavior of chemically treated and untreated hemp fiber reinforced composites was investigated.

The morphology of the composites was studied to understand the interaction between the filler and polymer. A series of dynamic mechanical tests were performed by varying the fiber loading and test frequencies over a range of testing temperatures. It was found that the storage modulus (E’) recorded above the glass transition temperature (Tg) decrease with increasing temperature. The loss modulus (E”) and damping peaks (Tan δ) values were found to be reduced with increasing matrix loading and temperature. Morphological changes and crystallinity of Composites were investigated using scanning electron microscope (SEM) and XRD techniques. The composites with Alkali and Benzoyl treated fibers has attributed enhanced DMA Results. In case of XRD studies, the composites with treated fibers with higher filler content show enhanced crystallinity.

Keywords: crystallinity, DMA, fiber treatments, hemp fiber, SEM, XRD.

1. Introduction Manufacturing high performance engineering materials from renewable resources is one ambitious goal currently being pursued by researchers across the world. The ecological benefits of renewable raw materials are clearly saved valuable resources are environmentally sound and do not cause health problems. Natural fibers have already established a track record as simple filler material in automobile parts. Natural fibers like sisal, hemp, kenaf, jute, coir, oil palm fiber have been proved to be good reinforcement in thermosets and thermoplastic matrices[1-4]. Previous studies have proved hemp fibers to be an effective reinforcement in polymer matrix[5].

The behavior of hybrid composites is a weighed sum of the individual components in which there is more favorable balance between the inherent advantages and disadvantages.

In an interesting study dynamic mechanical analysis of natural fibre based hybrid composites was performed and observed that the hybridization of nature fibre improved thermal and dynamic mechanical properties[6]. Glass/banana hybrid poly- ester composites are subjected to dynamic mechanical analysis over a range of temperature and three different frequencies[7]. Similar work on mechanical and dynamic mechanical properties of sisal/glass hybrid composites reported increase in storage and loss modulus with hybridization of sisal/polyester with glass fibres[8].

Mixing natural fibres like hemp and kenaf with thermoplastics put Flex Form Technologies[9] on the map and in the door panels of Chrysler’s Sebring convertible. In Germany, after authorization of hemp cultivation led to development of flax/hemp (50:50) needle felt for high-segment cars.

A landmark agreement between automobile giant Ford Automobiles supplier Visteon Automotive system and Kafue bio composites enhanced natural fibre composites

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applications in interior panels, linings and fittings. It is an important step towards higher performance of hybrid composites in automobiles applications. It is reported that presently, 27 components of a Mercedes S class are manufactured from natural fibre based composites with total weight of 43 kg[10]. The end of life vehicle[11] directive in Europe states that by 2015, vehicles must be constructed of 95% recyclable materials, with 85% recoverable through reuse or mechanical recycling[12]. The researchers[13-18] were carried out an improvement on dynamic mechanical properties of different polymeric composites system. The comparative studies in their investigation on the enhancement of damping in polymer composites have been suggested analyzing different fiber combinations. Extensive research work is being carried out regarding the visco elastic behavior of various polymer composites and blends[19,20].

Dynamic Mechanical Analysis (DMA) is a sensitive technique that characterizes the mechanical responses of materials by monitoring property changes with respect to the temperature and/or frequency of oscillation. The technique separates the dynamic response of materials into two distinct parts: an elastic part (E’) and a viscous or damping component (E”). The elastic process describes the energy stored in the system, while the viscous component describes the energy dissipated during the process. Mechanical loss factor (tan δ) is another useful parameter, which can be very useful[21-26].

The objective of this study was to investigate the influence of cellulose filler on dynamic mechanical behavior of Hemp/HFRCFE composites. HFRCFE composites using both (treated and untreated) Hemp fiber with cellulose filler is studied. Untreated HFRCFE was used in this study to obtain the elementary properties of Hemp-based HFRCFE with cellulose fillers.

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O O O O O O O O O O O O O O O O

Palanivel, A., Veerabathiran, A., Duruvasalu, R., Iyyanar, S., & Velumayil, R.

2. Materials and Methods 2.1 Materials Hemp fibre mats were purchased from Sheeba fibers and handicrafts, Poovancode, Tamilnadu, India. The epoxy resin employed in the present study is LY556 and the hardener HY951 was purchased from the Modern Scientific Pvt Ltd, Chennai, Tamilnadu, India. LY556 resin is a bi-functional epoxy resin i.e., diglycidyl ether of biphenyl-A (DGEBA) and HY951 is an aliphatic primary amine, viz., triethylene tetra mine – TETA with the mixing ratio is 10:1 w/w.Lyocell Powder (1.7 decitex) with around 12 µm was supplied by Simtek Lab Agencies, Navi Mumbai, India. Chemicals used for the surface modification of fiber are commercial Sodium Hydroxide and Benzoyl Chloride which were kindly supplied

by Simtek Lab agencies, Navi Mumbai, India.

Figure 1. Raw hemp fiber mat.

2.2 Treatment of fiber First the received Hemp fibers are washed with distilled water to remove the surface dirt present in the fibers and then the fibers are dried in an air circulating oven at a temperature of 1000C until it gains a fixed value of weight. Then the fibers are named as raw hemp fibers shown in Figure 1.

2.2.1 Bleaching treatment For this treatment 25g Hemp fibers were added to a 2 L solution containing 320 mL (30%; w/w) hydrogen peroxide and 1g sodium hydroxide and heated at 85 °C for 1 h[27,28].

During this process the fibers are cooked in the solution under gradual rise and fall of the temperature of the bath from 30 °C to 85 °C. This process of heating and cooling was done for a period of 1 h. Finally, the cooked fibers are removed from the mixture at a temperature of 30 °C. In order to remove excess mixture, the fibers are washed with distilled water. After washing, the fibers are again dried in an air circulating oven at a temperature of 100 °C until it gains constant weight[29]. Then the fibers are designated as bleached Hemp fibers.

2.2.2 Chemical treatments-alkaline treatment Alkaline treatment or mercerization is one of the most used chemical treatments of natural fibers when used to reinforce thermoplastics and thermosets. The important modification done by alkaline treatment is the disruption of hydrogen bonding in the network structure, thereby increasing surface roughness. This treatment removes a certain amount of lignin, wax and oils covering the external surface of the fiber cell wall, depolymerises cellulose and exposes the short length crystallites[13]. Addition of aqueous sodium hydroxide (NaOH) to natural fiber promotes the ionization of the hydroxyl group to the alkoxide[30]. A water-ethanol solution (80:20) is prepared (6% of NaOH) and stirred continuously for 1 hour. Later, fiber mats were immersed one by one in the solution. Finally after immersing all the fiber mats the object is left undisturbed for nearly 3 hours.

Then the fiber mats are washed several times with distilled water followed by drying it at 80°C for 5 hours in a hot air oven[31-34]. Then the fibers are designated as Treated Hemp fibers shown in Figure 2.

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Figure 2. Treated hemp fiber mat.

2.2.3 Benzoylation treatment Benzoylation is an important transformation in organic synthesis[35]. Benzoyl chloride is most often used in fiber treatment. Benzoyl chloride includes benzoyl (C6H5C=O) which is attributed to the decreased hydrophilic nature of the treated fiber and improved interaction with the hydrophobic PS matrix[35]. Benzoylation of fiber improves fiber matrix adhesion, thereby considerably increasing the strength of composite, decreasing its water absorption and improving its thermal stability[36,37]. It was observed that the thermal stability of treated composites were higher than that of untreated fiber composites[36]. The fiber was initially alkaline pre-treated in order to activate the hydroxyl groups of the cellulose and lignin in the fiber; then the fiber was suspended in 10% NaOH and benzoyl chloride solution for 15 min.

The isolated fibers were then soaked in ethanol for 1 h to remove the benzoyl chloride and finally was washed with water and dried in the oven at 80 °C for 24 h.

2.3 Composite fabrication The composite material is fabricated by using hand layup technique[38]. Composite fabrication using double weave and non-woven Hemp mats (150mm×150mm×1mm) was carried out in the square moulds of volume 350x350x3, 350x350x6 & 350x350x10 mm3. Initially, the mould was polished and mould releasing agent was applied on its surface. Resin, hardener mixture and Synthetic Cellulose Powder (10:1:4.3) Polímeros, 27(4), 309-319, 2017

Dynamic mechanical analysis and crystalline analysis of hemp fiber reinforced cellulose filled epoxy composite are spilled for every layer. Figure 3 shows the treated hemp fiber mat used for the material preparation.

Initially the fibres are dried in sun light to remove the moisture. The mould surface is cleaned and releasing agent (Wax) is applied. A thin layer of resin is also applied on the board. The woven roving NFRP are then completely filled with resin mixture, rolled to remove the entrapped air and to uniformly spread the mixture. In this way three layers of woven roving are placed one over the other to obtain top and bottom layers. A curing time of 24 Hrs at room temperature is given for the structures to obtain good strength. After curing for 24 Hrs, the required composite was obtained. By the same fabrication procedure, composites of different configurations by varying the ratio of cellulose powder and epoxy resin were fabricated. The same fabrication process was then carried out for fibers which were chemically treated. Thus, untreated and treated composites were prepared. Figure 4 shows the fabricated composite material.

Figure 3. Treated hemp fiber mat.

3. Testing of Composites All the tests were carried out as per ASTM Standards at Central Institute of Plastics Engineering & Technology, Chennai, India an ISO/IEC 17025:2005 – NABL Accredited Laboratory and ISO/IEC 17020 – NABCB Accredited Laboratory.

3.1 Dynamic Mechanical Analysis (DMA) A DMA is a precision instrument used to measure mechanical and visco elastic properties both on rigid and soft solid materials, under controlled temperature settings. Concretely, storage modulus and loss modulus together with temperature are key parameters in the Frequency Sweep-Temperature Step test, while stress and strain parameters characterize the Stress-Strain test. The samples are mounted on a clamp with a stationary and a movable part connected to the drive motor. Thus, the motor affects directly the deformation of the sample. The way it works is simple: the drive motor delivers force or stress to the tested sample while an optical encoder measures the resulting displacement. Furthermore, it includes an air bearing of nitrogen to assure a smooth and continuous delivery of force without noise. The instrument used to perform Dynamic Mechanical Analysis tests in this investigation is the TA Q800 DMA. To acquire and manage the results, this machine uses software which allows the user to visualize the obtained results, etc. To control the DMA, the user is able to choose between this software and the DMA’s touch-screen when configuring test parameters. Basic instrument characteristics are presented in Table 1 and those focused specifically on our tests in Table 2.

Dynamic mechanical analysis (DMA) of HFRCFE composites was performed with a single cantilever mode at constant amplitude of 15µm with 1% strain rate over a frequency of 0.1 Hz to 100Hz purging liquid nitrogen.

The heating rate was 5˚C/min. Before each measurement, the DMA instrument was calibrated to have the correct clamp position and clamp compliance. The specimen dimensions were 35 mm × 13 mm × 3 mm. The DMA experiments were carried out in the temperature range from Polímeros, 27(4), 309-319, 2017

Figure 4. Fabricated HFRCFE specimens.

Table 1. Basic instrument characteristics.

Operating environment conditions Temperature Range Displacement Range Loading

Temperature: 15-30 °C, Relative Humidity: 5-80% (non condensing) -145 to 600 ºC 25 mm 0.001 to 18 N

Table 2. Specific characteristics of rectangular samples.

Clamp Sample Length Sample Width Sample Thickness

3 Point Bending 35 mm (distance between support struts) 13 mm (varies slightly for each sample) 3 mm (varies slightly for each sample)

-30˚C to 200˚C under ASTM D4065 standards. The storage modulus (E’) & loss modulus (E”) and the mechanical loss factor (tan delta), as a function of temperature (T), were determined by dynamic mechanical analyzer were plotted versus temperature.

3.2 X-Ray Diffraction (XRD) study XRD measurements were made using Philips X’Pert powder diffraction system (Philips Analytical, The Netherlands) equipped with a vertical goniometer in the Bragg-Brentano focusing geometry. The X-ray generator was operated at 40 kV and 50 mA, using the CuKα line at 1.54056 Å as 311

Palanivel, A., Veerabathiran, A., Duruvasalu, R., Iyyanar, S., & Velumayil, R. the radiation source. Each powdered specimen was packed in a specimen holder made of glass. In setting up the specimen and apparatus, co planarity of the specimen surface with the specimen holder surface, and the setting of the specimen holder at the position of symmetric reflection geometry were ensured. The powders were passed through a 100 mesh sieve and were placed into the sample holder by the side drift technique[39]. The holder consisted of a central cavity.

In order to prepare a sample for analysis, a glass slide was clipped to the top face of the sample holder so as to form a wall. The powder sample was filled into the holder, gently tapped and used for XRD measurement. 10 mg of each sample was scanned at 25˚C from 10˚C to 70˚C (2θ) and in step size of 0.020 and count time of 2.00 s, using an automatic divergence slit assembly and a proportional detector. Relative intensities were read from the strip charts and corrected to fixed slit values. The Percentage crystallinity and crystallinity index was calculated as followed in (1).

I %Cr = 22 ×100 I 22 + I18

I 22 − I18 C.I . = (1) I 22

%Cr – Percentage of Crystallinity, C.I. – Crystallinity Index, I22 – Relative Intensity @ 220 & I18 - Relative Intensity @ 180.

3.3 Morphological study (Scanning Electron Microscope – SEM) To illustrate the effect of surface treatment of the fibre, the failure surfaces of the specimens subjected to test were analyzed using a JEOL scanning electron microscope (SEM).

In SEM a fine probe of electrons scans the surface of the sample and the signals emanating from the incident site are processed and quantized. All specimens were sputtered with 10 nm layer of gold prior to SEM observations. Each specimen was mounted on the aluminum holder of the microscope using double sided electrical conduction carbon adhesive tabs. The accelerating voltage of 5-15 kV was employed.

The SEM analyses of both raw and surface treated fibre reinforced composites were compared.

4.1 Dynamic mechanical analysis 4.1.1 Storage modulus (E’) The dynamic storage modulus (E’) is defined as the stress in phase with the strain in a sinusoidal shearing deformation divided by the strain[40]. The variation in the storage modulus as a function of temperature for the studied composites is given in Figure 5. As the temperature increases, E’ decreases for all composites and this can be attributed to the increase in the molecular mobility of the polymer chains[41]. A prominent increase in storage modulus of the matrix in the elastomeric region with the incorporation of fibers is expected due to an increase in stiffness of the matrix with the reinforcing effect imparted by the fibers.8 There is a clear increase in E” with filler loading and the maximum E” values were found for the composite with the optimum cellulose filler content of all samples (i.e., XHFRCFE). This may be associated with the strong filler/matrix interaction and the elastic modulus of the treated hemp fiber[8,40-42].The drop in the modulus (E’) on passing through the glass transition temperature is comparatively less for reinforced composites than for unreinforced resin. This can be attributed to the combination of the hydrodynamic effects of the fibers embedded in a visco elastic medium and to the mechanical restraint introduced by the filler at the high concentrations, which reduce the mobility and deformability of the matrix.

Other authors have also reported similar observations[43]. at higher temperatures any water molecules adhering on to the fiber will get evaporated making the fiber stiffer.

4. Results and Discussion The Table 3 shows the seven specimens having different compositions with three samples each were tested in DMA, XRD & SEM Analysis.

Figure 5. Effect of temperature on storage modulus.

Table 3. Composition of different specimen (Percentage of Weight).

S.No 1 2 3 4 5 6 7

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Specimen U HFRCFE V HFRCFE W HFRCFE X HFRCFE Y HFRCFE Z HFRCFE UX (Untreated) HFRCFE

Hemp (%) 8 8 8 8 8 8 8

Cellulose Powder (%) (1) 34.25 31.5 28.75 26 23.25 20.5 26

Epoxy resin (%) (2) 52.5 55 57.5 60 62.5 65 60

Hardener (%) (3)

Proportion 3:2:1

5.25 5.5 5.75 6 6.25 6.5 6

10:1:6.5 10:1:5.7 10:1:5 10:1:4.3 10:1:3.7 10:1:3.2 10:1:4.3

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Dynamic mechanical analysis and crystalline analysis of hemp fiber reinforced cellulose filled epoxy composite Air-dried cellulose is rather resistant to mechanical impact and a large amount of mechanical energy has to be spent in order to destroy the macroscopic and microscopic structure.

This ultimately contributes to the improved modulus of the composite at high temperatures.

4.1.2 Loss modulus (E”) The loss modulus (E”) is defined as the stress 90˚ out-of-phase with the strain divided by the strain. It is a measure of the energy dissipated as heat per cycle under deformation, i.e., the viscous response of the material[40]. Figures 6 show the variation in E” with the temperature for the different composites. From these figures, it is clear that the incorporation of treated fiber causes a broadening of the loss modulus peak.

This may be attributed to the inhibition of the relaxation process within the composites as a consequence of a higher number of chain segments upon filler addition[44]. There is an apparent shift in Tg toward higher temperatures on increasing the filler content and the overall filler loading. This is primarily attributed to the segmental immobilization of the matrix chain at the fiber surface[45]. The loss modulus in the transition region is also higher for composites with higher filler content, which may be due to an increase in internal friction, promoting energy dissipation. The high modulus glass fibers introduce constraints on the segmental mobility of the polymeric molecules at the relaxation temperatures, but, probably, there are also other factors that lead to energy dissipation[46]. It can be observed that, for higher filler content, the loss modulus curve spreads over a wider distribution and shows a higher peak. This effect can be a consequence of the inhibition/ restriction of the relaxation process of the chain segments in the composites or due to an increase in the rigidity of chain segments, increasing material heterogeneity[47]. A higher peak height may be associated to a poor interface[48]. However, in general, it may also be attributed to an increase in the mobility of the polymer chains[35]. The width of the relaxation curve is characterized by the β factor of the Kohlrausch-Williams- Watts equation (KWW). A low β value implies a wider distribution, whereas a β value close to the unity means a perfectly narrow relaxation spectrum[49,50]. An increase in free volume usually results in a decrease in the number of cross-linking sites of the polymeric matrix, because there is less cross- linked space between fiber and filler incorporated epoxy matrix. Also, a tendency toward a α-transition shift is noted for higher temperatures increasing the filler content. The increase in width of the loss modulus curve is taken to represent the presence of an increased range of order. The greater constraints on the amorphous phase could give rise to higher or broader glass transition behavior Figure shows the maximum peak width is found to be for the composites with 43% filler loading.

4.1.3 Damping parameters (tan δ) The ratio between the loss modulus (E”) and the storage modulus (E’) is called the mechanical loss factor, or tan δ.

Figure 7 delineates the effect of temperature on tan δ.

Improvement in interfacial bonding in composites occurs as observed by the lowering in tan δ values. The damping properties of the material give the balance between the elastic and viscous phases in a polymeric structure[51].

In composites, damping is influenced by the incorporation, Polímeros, 27(4), 309-319, 2017

Figure 6. Effect of temperature on loss modulus.

Figure 7. Effect of temperature on damping factor.

type and distribution of fibers, as well as the filler/matrix interaction and the void content[52-56]. All materials exhibit a relaxation process, which is associated with the glass-rubber transition of the matrix. It has been observed that as the temperature increases, the damping values pass through a maximum in the transition region and then decrease in the rubbery region. This relaxation process, denoted as α, involves the release of cooperative motions of chains between cross links. Below Tg, damping is low because, in this region, the chain segments are in the frozen state.

Hence, the deformations are primarily elastic and the molecular slips resulting in the viscous flow are low. Also, in the rubbery region, the molecular segments are quite free to move and hence damping is low and thus there is no resistance to flow[44,57]. However, in the transition region, the molecular chains begin to move and every time a frozen segment begins to move its excess energy is dissipated as heat. In fact, a frozen-in segment in the glassy state can store more energy for a given deformation than a rubbery segment, which can move freely. In a region where most of the chain segments take part in a cooperative motion under a given deformation, maximum damping will occur[42,44,58].

The position and height of the tan δ peak are indicative of the structure and properties of a particular composite material.

Generally, composites have considerably less damping in the transition region compared to neat resin because the fibers 313

Palanivel, A., Veerabathiran, A., Duruvasalu, R., Iyyanar, S., & Velumayil, R. and fillers carry a greater amount of the load and allow only a small part of it to strain the interface[45]. Therefore, energy dissipation will occur in the polymer matrix at the interface and a stronger interface allows less dissipation. This may be due to a restriction of the movement of the polymer molecules due to the incorporation of the stiff fibers[47]. As in the case of the E” curves, the tan δ curves of the composites shift towards higher values for higher overall and treated fiber content.

In general, this is indicative of a poor interface. Since a lower peak height indicates a good interfacial adhesion[59], according to the results shown in this study other factors contribute to the softening of the interface. Generally, composites containing less fiber content exhibit higher peak heights. One reason for this may be that there is less matrix by volume with higher filler content, and there is more energy at the interface because of the increase in the interfacial area. The matrix dissipates more energy than composites, because the fibers carry a greater amount of the load, dissipating a small part of it to strain the interface. The width of the tan δ peak of the composites also becomes broader upon fiber treatments. At lower fiber concentrations, the packing of fibers will become inefficient leading to matrix-rich regions and therefore the matrix is not restrained by sufficient fibers and highly localized strains will occur[47]. Higher value was obtained for the matrix more cross linking sites, as a consequence, a more heterogeneous system is obtained.

Figure 8. XRD plots of surface treated and raw XHFRCFE.

4.2 X-ray diffraction analysis X-ray diffraction studies were performed on X-ray diffractometer. It is evident from Table 4 that UXHFRCFE at 2θ scale gave peaks at 22.0 and 18.0 with relative intensity is 1370 and 819 respectively. Percentage crystallinity (%Cr) and crystallinity index (C.I.) of UXHFRCFE are 62.59 and 0.40 respectively whereas percentage crystallinity of Surface treated HFRCFE Composites were 71.59, 68.31,66.03, 64.52, 63.93 and 63.18 Whereas crystallinity index of Surface treated HFRCFE Composites are 0.6, 0.54, 0.49, 0.45, 0.44 and 0.42 respectively. The counter reading at peak intensity at 22˚ is said to represent the crystalline material and the peak intensity at 18˚ corresponds to the amorphous material in cellulose[60,61], Percentage Crystallinity[62] and crystalline index[63,64] were calculated and tabulated in Table 4. A poor crystallinity index in case of untreated hemp fibers reinforced in cellulose filled epoxy matrix means poor order of cellulose crystals to the fiber axis during treatments, indicated by the lower crystallinity index. Thus clearly indicate that cellulose crystals are better oriented in hemp fibers followed by alkali and benzoyl treated hemp fibers reinforced in cellulose filled epoxy composites.

The graph of raw and surface treated HFRCFE Composites is shown in Figures 8-10.

Figure 9. XRD plots of U, V & XHFRCFE composites.

Figure 10. XRD plots of W, Y & Z composites.

Table 4. XRD data of untreated and treated HFRCFE composites.

S.No

Specimen

I22 @ 2θ

I18 @ 2θ

%Cr

C.I

1 2 3 4 5 6 7

U HFRCFE V HFRCFE W HFRCFE X HFRCFE Y HFRCFE Z HFRCFE UX (Untreated) HFRCFE

1283 1351 1422 1625 1595 1515 1370

724 695 782 645 740 883 819

63.93 66.03 64.52 71.59 68.31 63.18 62.59

0.44 0.49 0.45 0.6 0.54 0.42 0.4

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Dynamic mechanical analysis and crystalline analysis of hemp fiber reinforced cellulose filled epoxy composite

Figure 11. SEM photographs of raw and treated HFRCFE composites.

X-ray results for treated composites which show an increase in the ‘crystallinity’ index indicate improvement in the order of the crystallites as the cell wall thickens upon surface treatments and increase in the filler content.

These treatments were reported to reduce the proportion of crystalline material present in plant fibers, and increase in filler loading as observed by several researchers[65].

The increase of crystallinity index indicated that the fillers in matrix induced the crystallinity and it increase due to the removal of amorphous materials like hemicellulose, lignin, and some other non-cellulosic material by surface treatment of fibers.

4.3 Morphological analysis (SEM Analysis) Scanning electron microscopic (SEM) provide an excellent technique for the study of surface morphology of cellulose filled epoxy matrix reinforced with treated hemp fibers. Surface morphologies of untreated and surface treated HFRCFE Composites are presented in Figure 11.

These results suggest that cellulose in raw Hemp fiber is held together by means of binding components such as lignin, pectin, etc, which are removed after treatments. On the other hand, treated fiber showed a comparatively smoother surface with narrow fiber-thickness. Surface morphologies of treated HFRCFEs were attributed to the removal of lignin, pectin and hemicelluloses and agree with our observed surface structures.

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The SEM Examinations of fractured surfaces of untreated fiber based composites with cellulose fillers are presented in Figure Composites with less percentage filler content revealed poor interfacial bonding. Micrographs indicated fiber pull out, debonding, delamination and fiber breakage.

On the other hand treated fiber based composites showed much less fiber pull out indicating the better interfacial adhesion between fiber and matrix due to treatment.

5. Conclusions In this article dynamic mechanical properties of hemp fiber reinforced cellulose filled epoxy hybrid composites are described. The effect of hybridization on the dynamic mechanical properties was studied in detail (Figure 12, 13).

As expected, the storage and loss modulus decreased with the increase in temperature, which is associated with a softening of the matrix at higher temperatures.

The storage modulus increased with increasing filler loading and this was due to the reinforcement effect imparted by the fillers which are more rigid than the matrix. The loss modulus curves were found to distribute over a wider range and reach higher peak values for higher filler content.

These curves, which are indicative of the dissipated energy, were found to be shifted to higher positions following the incorporation of more cellulose fillers into the composites.

The incorporation of treated hemp fibers also caused a broadening of the loss modulus peak which was attributed to 315

Palanivel, A., Veerabathiran, A., Duruvasalu, R., Iyyanar, S., & Velumayil, R.

Figure 12. DMA analysis of HFRCFE composites.

the inhibition of the relaxation process within the composites.

The higher loss modulus at the relaxation temperature was associated with an increase in internal friction which enhances energy dissipation. Additionally, despite the presence of high-modulus cellulose fillers, which introduce constraints on the segmental mobility of the polymeric molecules at the relaxation temperature, there are other factors that lead to energy dissipation. There was a shift in the glass transition towards higher temperatures on increasing the overall filler loading and the treated fiber content. This is because of restrictions imposed on the mobility of the polymer molecules at the interface. The highest activation energy values were obtained for the composites with higher glass fiber content.

Figure 13. DMA analysis of UXHFRCFE composites.

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It is believed that the increase in the crystallinity index obtained by X-ray diffraction is in actual fact an increase of PolĂ­meros, 27(4), 309-319, 2017

Dynamic mechanical analysis and crystalline analysis of hemp fiber reinforced cellulose filled epoxy composite the order of the crystallite packing rather than in increase in the intrinsic crystallinity. The surface of hemp fiber becomes rougher after treatments in comparison with smooth and clear surface of raw hemp fibers. The removal of surface impurities on plant fibers may be an advantage for fiber to matrix adhesion as it may facilitate both mechanical interlocking and the bonding reaction due to the exposure of the hydroxyl groups to chemicals such as resins and fillers. Better adhesion between filler and matrix, as evidenced by scanning electron microscopy, was found in the case of composites with treated fiber and filler content, while the composite materials without fiber treatments showed debonding and agglomeration of fiber, thereby decreasing their interfacial adhesion. Removal of non-cellulosic compounds is also suspected to increase the amount of OH groups exposed on the fiber surface, which could assist in mechanical enlargement with the matrix in the presence of cellulose filler, as evidenced by the increase in composites strength.

6. Abbreviations DMA – Dynamical Mechanical Analysis, SEM-Scanning Electron Microscope, XRD – X Ray Diffraction, HFRCFE-Hemp Fiber Reinforced Cellulose Filled Epoxy Composites.

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Received: Apr. 20, 2016 Revised: June 29, 2016 Accepted: July 12, 2016

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http://dx.doi.org/10.1590/0104-1428.05316

O O O O O O O O O O O O O O O O

Production of biodegradable starch nanocomposites using cellulose nanocrystals extracted from coconut fibers Jamile Costa Cerqueira1, Josenai da Silva Penha1, Roseane Santos Oliveira1, Lilian Lefol Nani Guarieiro2, Pollyana da Silva Melo2, Josiane Dantas Viana3 and Bruna Aparecida Souza Machado1* Applied Research Laboratory of Food and Biotechnology, Department of Food and Beverages, Centro Universitário – SENAI-CIMATEC, Salvador, BA, Brazil 2 Integrated Laboratory of Applied Research in Chemistry, Department of Automotive, Centro Universitário – SENAI-CIMATEC, Salvador, BA Brazil 3 Materials Laboratory, Department of Materials, Centro Universitário – SENAI-CIMATEC, Salvador, BA, Brazil

1

*brunam@fieb.org.br

Abstract Different polymeric matrices have been investigated for use in the development of biodegradable films. The incorporation of cellulose nanocrystals in such films has particularly attracted attention because of the potential for achieving improved properties of starch nanocomposites. In the present study, cellulose nanocrystals were extracted from coconut fibers and incorporated in cassava and potato starch films at different concentrations. The properties of the different nanobiocomposite films were comparatively evaluated, including their barrier and mechanical properties. All the films, regardless of the nanocrystal concentration, were found to exhibit low solubility in water, with increased moisture content particularly observed in the films with higher nanocrystal concentrations. The potato starch film with the lowest nanocrystal concentration was found to exhibit the best mechanical properties. The observations of this study indicated that the source of the starch and the nanocrystal concentration determined the properties of the nanobiocomposite films.

Keywords: lignocellulosic fiber, biodegradable film, polymeric matrix, nanocrystal.

1. Introduction Plastics are the most commonly used materials in the food industry, owing to their moldability, manageability, low cost, and suitable mechanical and chemical properties[1].

However, the diversity of resins used for their production poses challenges to their separation for reuse[2]. This has contributed to the environmental problems associated with discarded plastic packaging, with the significant buildup of residues requiring about 100-450 years for their natural decomposition[3]. The increasing demand for environmental preservation has thus necessitated research into alternatives to conventional plastics, such as varieties derived from petroleum derivatives using biodegradable and ecologically friendly materials.

The technological adaption of biomaterials, particularly polymers, for the development of flexible films has been widely studied in recent years[4–6]. Several biopolymers such as polysaccharides, proteins, and lipids have been utilized as polymeric matrixes for the development of alternative biodegradable packaging, owing to their availability, renewability, low cost, environmental friendliness, and biodegradability[7-9]. Among such agriculturally sourced materials, starch is considered to be favorable for the development of biocomposites[10,11]. Starch films are particularly characterized by transparency, non-toxicity, and low cost. However, they also have some limitations, the most prominent of which are low flexibility, high permeability to water vapor[2,12], and inadequate mechanical and barrier resistance. This causes them to fall short of market

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expectations, especially in the food industry. Consequently, several studies have been conducted on the use of reinforcing additives such as fibers[13,14] and cellulose nanocrystals[15-17] to improve the mechanical, thermal, and barrier properties of starch-based films.

Cellulose is the most important polymer and is used in various fields, especially because of its availability, biocompatibility, and biodegradability[18,19]. Cellulose fibers have an inherent structural hierarchy originating from their differing biological sources. The production of cellulose nanocrystals from different lignocellulosic fibers has attracted interest in the scientific and industrial community in recent years. This has been owing to their mechanical and thermal properties, cost effectiveness, and versatility[20,21]. Nanocrystals are generally characterized by high solidity and crystallinity, which afford sustainability.

When incorporated in films with renewable polymeric matrixes, they produce total biodegradable composites[22] referred to as nanobiocomposites. As with any composite material, the properties of the nanobiocomposites depend on the properties of the individual components (matrix and reinforcement), the composition (volumetric fractions of the components), the morphology of each constituent phase (spatial arrangement, dimensions, and crystallinity), and the interfacial properties[23].

It is important to highlight that the development of new composites with natural fiber reinforcement is quite promising, bearing in mind that, in addition to the renewability of the

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Production of biodegradable starch nanocomposites using cellulose nanocrystals extracted from coconut fibers fiber sources, they also cost less than synthetic fibers[24,25].

Incidentally, Brazil is acknowledged as one of the largest agricultural producers in the world, contributing a large part of the millions of tons of lignocellulosic biomass produced globally from different sources[23]. The green coconut bark is an important source of natural fiber, and its utilization for the production of cellulose nanocrystals offers the benefit of reduced disposal of biomass[25-27].

Further, it should be noted that the use of biodegradable packaging in the food industry promises to enhance the preservation of the packaged products. It would also lead to the creation of new consumer markets as the differentiated biodegradable properties of the films fit new consumer profiles, especially those who are prepared to pay more for products or services with a green stamp (a certificate granted to products manufactured for sustainable development)[28,29].

It is noteworthy that recent studies have highlighted the benefits of incorporating cellulose nanocrystals in flexible films produced from cassava starch, namely, the improved properties afforded by the produced biodegradable starch nanocomposites[5,6,10,12]. In the present study, we incorporated cellulose nanocrystals produced from green coconut fiber in flexible films with polymeric matrixes of tapioca starch and potato starch, respectively. We evaluated the effects of the starch source and the concentration of the cellulose nanocrystals on the properties of the films, and determined the best matrix and nanocrystal concentration. Glycerol was used as plasticizer in the development of the flexible films.

The objective of the study was to characterize green coconut fiber cellulose nanocrystals with respect to their size.

2. Materials and Methods 2.1 Extraction of cellulose from green coconut fiber The extraction of the cellulose pulp from the green coconut fiber was performed based on the works of Rosa et al.[26], Samir et al. [30], and Machado et al.[5]. The coconut fibers were dried in a greenhouse and ground in a liquidizer to obtain a fine particulate. The particulate was further dried and 30 g of it was washed by constant agitation for 4 h in 1.2 L of a solution of 2% NaOH at 80°C. The fine particulate was then filtered and the washing process was repeated three more times to completely remove the water-soluble agents.

The filtered particulate was subsequently further washed in distilled water to obtain the cellulose pulp. This was followed by delignification (or whitening) of the pulp using a mixture of 0.3 L of 1.7% sodium hypochlorite and 0.3 L of a buffer solution. The cellulose pulp was constantly agitated in the mixture for 6 h at 80°C, and then filtered and dried in a greenhouse. It was finally pulverized in a mill (Cadense Ltd, Brazil).

2.2 Preparation of cellulose nanocrystals The cellulose nanocrystals were prepared by acid hydrolysis using 64% H2SO4[5,6,26], with 12 mL of cellulose pulp per gram of the acid subjected to constant agitation for 5 min at 50°C. After acid hydrolysis, the samples were filtered and their concentrations adjusted. The samples were then centrifuged for 10 min at 4400 rpm and 10°C to separate the crystals. This procedure was repeated eight times, when Polímeros, 27(4), 320-329, 2017

no more supernatant was produced. The suspensions were then dialyzed using the D9777-100 FTO cellulose membrane with 12.000 Da cut off (Sigma-Aldrich) until their pH was approximately 7, after which the samples were placed in an ultrasonic bath for 20 min.

2.3 Preparation of nanobiocomposites The nanobiocomposites were prepared by the casting method, which is used for the preparation and drying of films. The process involved the preparation of a film-forming solution in which the polymeric matrix (4%, g/100 g)— manioc starch (Cargill Agrícola SA) or potato starch—and glycerol (Synth, 1.0%, g/100 g) were initially dissolved.

The solution was prepared with distilled water. The mixture was then added to a suspension of the cellulose nanocrystals containing varying amounts of the nanocrystals (0.5%, 1.0%, and 1.5% (g/100 g)) (Table 1). The new mixtures were then heated to the gelatinization temperature of the starch (~70°C) under constant manual agitation. 45 g of the mixtures was then weighed in polystyrene Petri dishes and dehydrated in a greenhouse under airflow (35±2°C) for 18-20 h. Other films were produced without the addition of cellulose nanocrystals for use as controls. Before their characterization, the prepared nanobiocomposites were stored at 60% humidity and 23°C for 10 days in a desiccator containing a saturated solution of sodium chloride.

2.4 Characterization of coconut fibers and cellulose nanocrystals The main components of the fibers obtained from green coconut bark (lignin, hemicellulose, and cellulose) were characterized using the ANKOM A200 Fiber Analyzer.

The methodology proposed by Van Soest et al.[31] and Goering and Van Soest[32] was employed. The fibers and cellulose were also characterized using a scanning electron microscope (SEM) (model JEOL JSM-6390LV). The fibers and cellulose samples used for the SEM characterization were metalized with gold in a “sputter coater” (model SCD 50, Balzers) using argon plasma, and then left to rest for 24 h.

To determine the birefringence of the suspensions of the cellulose nanocrystals in water, two films of crossed polarizers were used. The polarizer films were positioned perpendicularly to each other and a light was directed at one of them. The suspension sample was then interpolated between the two films[33].

Table 1. Design of the experimental nanobiocomposites.

Formulation FM05 FM10 FM15 FB05 FB10 FB15 Control FM Control FB

Manioc starch

Potato starch

Cellulose nanocrystals

(%, g/100 g) (%, g/100 g) (%, g/100 g) 4.0 0.5 4.0 1.0 4.0 1.5 4.0 0.5 4.0 1.0 4.0 1.5 4.0 4.0 -

Glycerol (%, g/100 g) 1.0 1.0 1.0 1.0 1.0 1.0 1.0 1.0

321

Cerqueira, J. C., Penha, J. S., Oliveira, R. S., Guarieiro, L. L. N., Melo, P. S., Viana, J. D., & Machado, B. A. S. The suspension of the coconut cellulose nanocrystals was examined by an SEM to determine the crystal length (L), diameter (D), and aspect ratio (L/D), and the state of the aggregation of the nanocrystals. The nanocrystals were visualized by coloring with 2% uranyl acetate (UA).

The coloring was done by mixing the suspension of the cellulose nanocrystals with an equal volume of UA. 10 µL of the mixture was then distributed through 400‑mesh copper grids and left to stand for 30-60 s. The grids were subsequently dried and visualized by a transmission electron microscope (TEM) (BX-51, Olympus) operated in the clear field mode at 80 kV. The length and width of the nanocrystals were directly measured from the transmission electron micrograph. The mean and standard deviation for 30 measurements were determined.

2.5 Characterization of nanobiocomposites and controls The thicknesses of each pre-packaged nanobiocomposite film (60% humidity, 25°C) were determined by calculating the average thickness for six measurements at random points.

A flat-head digital micrometer (Mitutoyo) with a resolution of 1 µm was used for the measurements. The water activity (aw) of the nanobiocomposite film was also measured using a Lab Master decagon with an electrolytic measurement cell (CM-2, TECNAL), with the temperature controlled to 25°C[34,35]. Further, the humidity and total solid content of the nanobiocomposite film were determined by infrared drying using a Shimadzu infrared scale, with the intensity of the radiation adjusted to raise the temperature of the sample to 105°C. The water solubilities of the different formulations were determined using the method proposed by Gontard et al.[36].

The traction assays were performed using a texture analyzer (model CT310K, Brookfield) with a maximum load of 10 kN. The ASTM D-882 standard that was adopted for the assays stipulates a loading speed of 0.50 mm/s, temperature of 25°C, trigger load of 7 g, test probe tip of TA3/100, and the use of a TA/TPB probing device. The traction assays were performed using six proof bodies for each sample.

The proof bodies were of length 80 mm and width 25 mm, and their strength was calculated by dividing the maximum applied force by the area of the film (width × thickness).

The strain at breakage was calculated by dividing the final length by the projection of the probe tip (50 mm) and multiplying by 100[37].

2.6 Statistical analysis The results of the above investigations were expressed in the form of mean ± standard deviation (n = 3). Statistica 6.0 (StatSoft, Tulsa) was used for the statistical analyses of the results. Analysis of variance (ANOVA) and Tukey’s tests were used to determine the significant differences between means (p < 0.05).

3. Results and Discussions 3.1 Characterization of coconut fibers and nanocrystals Vegetable fibers are mainly composed of cellulose, hemicellulose, and lignin, together with other smaller constituents. The physical and chemical properties of 322

a particular fiber depend on its source and processing.

The analysis of the chemical composition of the green coconut fiber in this study revealed cellulose, lignin, and hemicellulose contents of 32.0%, 38.0%, and 0.25%, respectively (Table 2). Trugilho et al.[38] reported that the lignin and cellulose contents of young plants are variable, and only stabilize with maturation. Corradini et al.[39] particularly found that the lignin and cellulose contents of a young green coconut fiber varied within 37.2%-43.9% and 31.5%-37.4%, respectively. Similar results were obtained

by Rosa et al.[26,40], who highlighted the high lignin content of green coconut fiber (Table 2) and the potential for its use in the reinforcement of polymeric materials[39].

Figure 1 shows SEM images of green coconut fiber (Figure 1A) and the extracted cellulose (Figure 1B).

The micrographs of the in natura fiber (Figure 1A) reveal rugged surfaces covered by layers of wax and extracts, as well as amorphous constituents such as lignin and hemicellulose. The micrographs of the whitened cellulose (Figure 1B) reveal the presence of cellulose fibrils and provide evidence of disaggregated ruggedness, which was due to the conversion of cellulosic elements to fibrils by the removal of poliosis and lignin during the whitening process. Pores and/or orifices can also be observed on the rugged surface. As also evidenced by Alemdar and Sain[18], these images suggest that the chemical treatment partially removed impurities such as hemicellulose and lignin, which are the “compacting” components of the fibers. The efficient removal of lignin by the whitening of the fibers is extremely important to obtaining nanocrystals by hydrolysis.

The visualization of the suspension of the cellulose nanocrystals obtained from the green coconut fiber (concentration of 0.06 g/10 mL) through polarizers revealed a nematic phase, which was directly produced by the light birefringence and is considered to confirm the presence of nanocrystals (Figure 2).

Mesquita et al.[41] and Alves et al.[17] similarly used crossed polarizers to visualize the birefringence phenomenon in a 4% suspension of cellulose nanocrystals obtained from eucalyptus. Birefringence (or optical anisotropy) is a property of transparent crystals that enables them to decompose a light ray into two crossed polarized rays. Observation of the birefringence in a cellulose nanocrystal suspension is thus an effective means of evaluating the nanocrystal formation[42], and hence assessing the efficiency of the hydrolysis process.

It has been observed that suspended cellulose nanocrystals exhibit a tendency to align themselves, probably due to their high solidity and length-to-diameter ratio[30]. It is also very important to determine whether the cellulose nanocrystals are well dispersed in suspension, because this is required to achieve good results when they are incorporated in a polymeric matrix for mechanical reinforcement[43,44].

Figure 3 shows the TEM micrographs of the cellulose nanocrystals, from which the length (L), diameter (D), Table 2. Chemical composition of green coconut fiber.

Cellulose (%) 32.0 32.0 37.0 31.5%-37.4%

Lignin (%) 38.0 40.0 32.5 37.2%-43.9%

Reference This study Rosa et al.[40] Rosa et al.[26] Corradini et al.[39]

Polímeros, 27(4), 320-329, 2017

Production of biodegradable starch nanocomposites using cellulose nanocrystals extracted from coconut fibers

Figure 1. SEM images of the (A) in natura green coconut fiber and (B) extracted cellulose (B) (Ranges of 1 mm and 100 µm).

Figure 2. Illustration of the suspension of cellulose nanocrystals obtained from green coconut fiber: (A) positions of the Polaroid and sample; (B) sample agitation by birefringence evidencing the presence of the nanocrystals; and (C) magnification of the image in (B) (the arrows indicate the locations of the birefringence phenomena).

Figure 3. TEM images of the nanocrystals obtained from green coconut fiber cellulose, showing (A) agglomeration of the nanocrystals and (B) some isolates (Range: 200 nm).

Polímeros, 27(4), 320-329, 2017

323

Cerqueira, J. C., Penha, J. S., Oliveira, R. S., Guarieiro, L. L. N., Melo, P. S., Viana, J. D., & Machado, B. A. S. and aspect ratio (L/D) of the nanocrystals and their state of aggregation were determined. The images confirm the effectiveness of the acid hydrolysis used to produce the cellulose nanocrystals from the green coconut fiber.

The images indicate that the aqueous suspensions contained needle-like nanocrystals, consisting primarily of individual fibrils and their aggregates. The cellulose nanocrystals were formed by nanoparticles containing several hydroxyl groups on their surfaces and with a large specific area. Aggregation of the nanocrystals was thus common, as expected, mainly due to the strong hydrogen links formed between cellulose nanoparticles[45].

The coconut cellulose nanocrystals examined in this study had length L values of 89–320 nm (average of 264.9±23.0 nm) and average diameter D of 8.10±1.21 nm.

These results agree well with previously obtained data for cellulose nanocrystals extracted from coconut fibers, namely, L values of 80–500 nm and D values of 4-9 nm[5,26,40].

Silva et al.[45] also reported an L value of 145 nm and average D of 1.5 nm for nanocrystals obtained from eucalyptus, while Costa et al.[6] reported L = 157 nm and average D = 5.7 nm for nanocrystals obtained from licuri fiber. The average aspect ratio (L/D) determined in the present study was 32.7±5.1, which confirms the potential of the coconut cellulose nanocrystals for use as reinforcement agents for polymeric matrixes in the development of composites, as proposed by Rosa et al.[40] and Machado et al.[5], who reported L/D values of 39±16 and 38.9±4.7, respectively.

control films, as well as among their respective samples.

The potato starch films were also found to be thicker than the manioc starch films (see Table 3). The mechanism of the film formation is known to be dependent on the concentration of the solids in the formulation and the amylose content.

The potato starch formulation appeared to be more viscous than that of the manioc starch, suggesting a higher amylose concentration of the former, and hence the thicker films.

With regard to the formulations containing the same matrix, namely, the same starch content and plasticizer concentration, variation of the drying method (casting) directly resulted in varying film thicknesses, as previously reported[6,47-49].

Fakhouri et al.[37] obtained varying film thicknesses between 0.039 and 0.081 mm for starch films plasticized with gelatin, while Reis et al.[14] obtained thicknesses of 0.11-0.12 mm for manioc starch films containing mango pulp and mate extract.

The water solubility of a film is an important property that determines its applicability for food packaging[50].

In some cases, total water solubility could be beneficial, such as in the case of semifinished food products that require further cooking. However, when the food item is liquid or aqueous, highly soluble polymeric films are undesirable[51].

The percentage solubility of the nanocomposites developed

3.2 Characterization of nanobiocomposites The films produced from the six formulations of nanobiocomposites prepared in this study using different starch types (manioc and potato), as well as the control films, were examined to determine their barrier properties (thickness, water activity, water solubility, and humidity) and mechanical properties (strength and strain at breakage).

An attempt was also made to evaluate how these parameters were affected by the addition of the coconut fiber cellulose nanocrystals. The produced starch films were found to be homogeneous, transparent, and visually attractive. Figure 4 shows an example of the manioc starch film with 0.5% nanocrystals (FM05). The determined barrier and mechanical properties of the different films are given in Tables 3 and 4.

Control of the thickness is extremely important to maintaining uniformity of the films and comparing their properties[46]. Significant difference was observed between the average thicknesses of the FM10 (0.08 mm) and FB (0.12 mm)

Figure 4. Nanobiocomposites obtained by the incorporation of 0.5% coconut cellulose nanocrystals in a manioc starch matrix plasticized with glycerol (FM05).

Table 3. Barrier properties (average ± standard deviation) of the films produced using different nanobiocomposite formulations.

Formulation Control FM FM05 FM10 FM15 Control FB FB05 FB10 FB15

EP 0.08±0.01ª 0.08±0.01ª 0.08±0.01b 0.09±0.02c 0.12±0.02d 0.09±0.01c 0.10±0.02e 0.10±0.01e

aw 0.65±0.01ª 0.58±0.01b 0.68±0.01c 0.71±0.01c 0.69±0.01c 0.66±0.02c 0.70±0.01c 0.68±0.02c

UM 15.31±0.78a 16.21±0.97b.c 18.33±1.29b 17.98±0.26b 16.60± 0.50c 18.87±0.61b 18.26±0.91b 18.46±0.38b

ST 84.68±0.78a 83.79±0.97b.c 81.67±1.29b 82.01±0.26b 83.39±0.50c 81.12±0.61b 81.74±0.91b 81.53±0.38b

SL 30.55±0.83a 30.54±0.83a 31.56±0.80b 32.16±0.65c 28.30±0.78d 29.52±0.46d 25.30±0.36e 29.17±0.65d

EP - thickness (mm), aw - water activity (%), UM - humidity (%), ST - total solids (%), SL -solubility (%). No significant difference between values with the same superscript letter in a column (p > 0.05), according to Tukey’s test with 95% confidence.

324

Polímeros, 27(4), 320-329, 2017

Production of biodegradable starch nanocomposites using cellulose nanocrystals extracted from coconut fibers Table 4. Mechanical properties (average ± standard deviation) of the films produced using different nanobiocomposite formulations.

Formulation

ST

↑ST

ε

↓ε

Control FM

3.23±0.51a

---

26.39±5.04b

---

FM05

3.65±0.68b

13.00

23.90±1.67c

9.43

FM10

4.41±0.30c

36.53

17.54±0.80d

33.53

FM15

3.80±0.94b

17.64

24.71±3.18bc

6.36

Control FB

4.07±0.20

---

30.81±4.44ª

---

FB05

8.20±0.52e

101.47

23.71±3.54c

23.04

FB10

5.43±0.24f

33.41

26.50±3.75b

13.98

FB15

4.09±0.64

0.31

30.40±1.77ª

1.33

d

d

ST - strength (MPa), ↑ST - increase in strength relative to control (%), ε - strain at breakage (%), ↓ε - decrease in strain at breakage relative to control (%). No significant difference between values with the same superscript letter in a column (p > 0.05), according to Tukey’s test with 95% confidence.

in the present study varied between 25.30 (FB10) and 32.16 (FM15), and may thus be considered as low-solubility films[52]. This indicates that they can be used to package a wide variety of food items, provided that their overall performance satisfies relevant legislation.

It was not possible to determine a direct effect of the nanocrystals on the water solubility of the present films.

However, the films containing the nanocrystals generally exhibited higher water solubility compared to the control films, although a lower solubility was observed for the film produced from the potato starch matrix containing 10% nanocrystals. The present solubility results are similar to those obtained by Jiang et al.[50] for films produced from pea starch reinforced with potato starch nanoparticles.

Considering the observations of Rubentheren et al.[53] and Pagno et al.[54] all the films containing cellulose nanocrystals were expected to exhibit significantly lower solubility. This is based on the stronger resistance of the films compared to the control films (Table 4), mainly due to the interaction between the nanocrystals and the starch chains. However, only the film produced using formulation FB10 exhibited such decreased solubility.

With regard to the water activity (aw) of the films, the values were found to vary between 0.58 (FM05) and 0.71 (FM15), with significant differences observed for only the films produced from the manioc starch formulations.

The aw was not found to be directly affected by the nanocrystal content of the films. The property is, however, an important consideration for food packaging[55] because the water content affects the growth of microorganisms. The property is thus directly related to the quality conservation of the packaged product. Silva et al.[55] found that starch films containing cellulose nanocrystals had aw values of 0.46-0.63, which substantiates the present observation that formulations containing higher percentages of cellulose nanocrystals produce films with lower aw values. Similar results were obtained by Costa et al.[6], namely, aw values of 0.44-0.49 for manioc films containing propolis extracts and reinforced with licuri cellulose nanocrystals.

All the films of the present study containing coconut cellulose nanocrystals were found to exhibit significantly higher humidity contents compared to the control films.

The humidity of the films produced from the control formulation FM was 15.31%, while those of the films Polímeros, 27(4), 320-329, 2017

produced from the manioc starch formulations containing cellulose nanocrystals varied between 16.21% and 18.33%.

Similar humidity values were obtained for the films produced from the potato starch formulations containing cellulose nanocrystals. The humidity of the films produced from the control formulation FB was 16.60%, while those of the films produced from the other formulations varied between 18.26 and 18.87%. These results indicate that the addition of coconut cellulose nanocrystals dispersed in water increased the humidity of the films. Considering the similar preparation conditions (amount of polymeric matrix, amount of plasticizer, and duration of drying), the higher humidity of the films containing cellulose nanocrystals may be due to the storage for ten days in a desiccator containing saturated sodium chloride solution before the characterization.

The films containing the cellulose nanocrystals might have absorbed more water in the desiccator than the control films.

Another possible explanation is that the nanocrystals that were incorporated into the plasticized polymeric matrix were dispersed in water, which could have increased the humidity of the films compared to the control films.

Expectedly, the strength and strain at breakage of the starch films were increased by the addition of the cellulose nanocrystals, as indicated in Table 4 and also previously reported[6,56,57]. The implied enhanced resistance to rupture is desirable and broadens the potential applications of the films to packaging and coating [17]. The strengths of the films ranged between 3.23 MPa (control FM) and 8.20 MPa (FB05). An increase of the cellulose nanocrystal concentration was not observed to increase the rupture resistance of the films, indicating that the addition of a large amount of nanoparticles could be counterproductive through their aggregation[58]. For the manioc starch films, the optimal nanocrystal concentration for mechanical reinforcement was determined to be 1.0%, while it was 0.5% for the potato starch films. In the case of the control films, those produced from potato starch were stronger. Overall, the potato starch films containing cellulose nanocrystals were generally the strongest. The mechanical properties of starch are known to be significantly affected by their botanical origin, specifically the native amylose and amylopectin contents[59]. The present observed film strengths could thus be explained by the fact that the molecular interactions among the starch chains, plasticizer, and dispersed nanocrystals in the potato starch films were more efficient. This may be due to the better 325

Cerqueira, J. C., Penha, J. S., Oliveira, R. S., Guarieiro, L. L. N., Melo, P. S., Viana, J. D., & Machado, B. A. S. dispersion of the nanocrystals within the matrix and stronger interfacial adhesion among the components of the complex system, which included the starch, glycerol, and nanocrystals.

Similar strength results were obtained by Alves et al.[17], who investigated the incorporation of eucalyptus nanocrystals in maze starch matrixes containing gelatin. They observed better dispersion of the nanocrystals in films with higher gelatin contents, resulting in better mechanical properties.

Oun et al.[60] also observed higher strength values of 22.7-33.7 MPa for agar-based nanocomposites and cotton cellulose nanocomposites. They found that the film traction increased with increasing nanocrystal concentration from 1% to 5%, but decreased significantly (p < 0.05) for 10% nanocrystals. Similar reports have been made by other researchers regarding the tensile properties of films with different matrixes containing cellulose nanocrystals[61-66].

The determined values of the strain at breakage ranged between 17.54% (FM10) and 30.81% (control FB). It was observed that the two strongest film types (FB05 and FM10) also exhibited the lowest breakage strains. This may be attributed to the higher rigidity of the films due to the strong interactions among the nanocrystals, plasticizer, and polymeric matrix. A small strain at breakage has actually been reported to be an indication of good interaction between the starch and nanocrystals[54,67,68].

4. Conclusions The combination of the availability of lignocellulosic fibers and the need for renewable resources for the production of polymers affords strong opportunities for adding value to agricultural products through technological innovation.

The results of this study, as well as those of previous works, specifically confirm the viability of developing biodegradable polymeric packaging with adequate characteristics.

The nanobiocomposite biodegradable films developed in the present study exhibited desirable barrier and mechanical properties including good transparency, manageability, homogeneity, adequate solubility, and high strength.

The addition of cellulose nanocrystals obtained from green coconut fiber to the polymeric matrixes particularly afforded efficient mechanical reinforcement. A low or moderate nanocrystal concentration was found to be more effective depending on the type of polymeric matrix (potato or manioc starch), with higher concentrations inducing undesirable agglomeration of the nanocrystals. The best barrier and mechanic properties were exhibited by the films produced using potato starch, attributable to the better interactions among the polymeric matrix, the cellulose nanocrystals, and the employed plasticizer. It should be noted, though, that the optimal composition of the formulation used to develop the film would depend on the intended use and application technique.

5. Acknowledgements The authors greatly appreciate the funding of this study by the National Service of Industrial Learning (Serviço Nacional de Aprendizagem Industrial) of the Regional Department of Bahia, the Foundation for Support to Research in Bahia (Fundação de Amparo a Pesquisa do 326

Estado da Bahia (FAPESB)), and the National Council of Technological and Scientific Development (CNPq).

They also thank the Gonçalo Moniz Research Centre – Fiocruz (Bahia) for their assistance with the microscopy analyses.

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Received: Apr. 22, 2016 Revised: Dec. 03, 2016 Accepted: Dec. 12, 2016

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http://dx.doi.org/10.1590/0104-1428.15616

O O O O O O O O O O O O O O O O

Layer-by-Layer technique employed to construct multitask interfaces in polymer composites Luísa Sá Vitorino1* and Rodrigo Lambert Oréfice1 1

Department of Metallurgical and Materials Engineering, Universidade Federal de Minas Gerais – UFMG, Belo Horizonte, MG, Brazil *luisavitorino@hotmail.com

Abstract The properties of glass fiber-reinforced polymer composites are closely related to the fiber-matrix interface. Interfacial treatments to improve mechanical properties are usually limited to enhance interfacial adhesion. In this work, Layer-by-Layer (LbL) technique was introduced to build a novel interface in polymer composites. Different numbers of bilayers of poly(diallyldimethylammonium chloride) and poly(sodium 4-styrenesulfonate) with carbon nanotubes were deposited through LbL on the surface of woven glass fibers (GFs). Polypropylene composites containing the modified GFs were prepared by compression molding. Thermogravimetric analysis, scanning electron microscopy and Raman spectroscopy proved that multilayers of polymers with carbon nanotubes could be deposited on GFs surface. Mechanical tests on composites with modified GFs revealed an increase in Flexural Modulus and toughness. The overall results attested that the LbL technique can be used to design interfaces with different compositions to perform diverse tasks, such as to improve the stiffness of composites and to encapsulate active nanocomponents.

Keywords: carbon nanotubes, composites, glass fibers, interface, layer-by-layer.

1. Introduction Polymer composites have intrinsic interfaces among their contituents that are responsible for many of the properties of the system. High levels of interfacial interactions between fibers and polymer matrices are considered critical for an efficient stress transfer and, consequently, high mechanical properties[1,2]. Many treatments have been studied and used to improve adhesion between the components of polymer composites, including the use of silane coupling agents, polymer grafts and plasma treatments[2-4]. These treatments are often successful in improving the adhesion between fibers and polymer matrices, but seldom provide other functionalities to the interface of polymer composites.

Glass fiber-reinforced polymers have been extensively used as materials for weight reduction with many engineering applications due to their high strength and specific stiffness.

Glass fibers (GFs) have developed an important role as reinforcement on components manufactured by diverse industrial sectors. Their high mechanical properties associated with good heat resistance and low cost have attracted glass fibers to attention[5-6]. Regarding structural applications (e.g. automotive, aerospace, construction), woven glass fiber fabric composites have been considered due to their deformation characteristics and stability. Moreover, polypropylene-glass fiber composites have been increasingly employed in advanced applications since polypropylene presents low density, low cost and easy processing[7].

When glass fibers are incorporated into a polymeric matrix, the properties of the fiber-matrix interface affect the composite performance significantly. A good interfacial adhesion can create an efficient load transfer from the polymer to the fiber, reducing stress concentrations and improving composite mechanical properties[1]. Glass fiber

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reinforced polymers began to present great improvements in their properties since the starting of using glass fibers with chemical treatments. Traditional treatments as sizings act by increasing the adhesion on interface but they are limited to increase some of the mechanical properties[5,8].

In this present work, we introduce a new way for modifying glass fibers by using the Layer-by-layer technique. Layer-by-Layer (LbL) technique is based on the consecutive adsorption of oppositely charged polyelectrolytes.

Two oppositely charged layers represent one bilayer which can be deposited until the quantity of interest. Compared to other self-assembly methods, LbL can be used to combine a variety of species in nanoscale, merging properties of each material to introduce specific new functions. The ability of LbL to control the thickness of the coating, the economic use of materials and the use of multiple nanocomponents represent a great advantage over other techniques[9-11].

In this work, the hypothesis that the LbL technique could be used to build designed interfaces in polymer composites was tested. The surface of glass fiber fabrics was modified by depositing, via the LbL technique, multilayers of poly(diallyldimethylammonium chloride) (PDDA) and poly(sodium 4-styrenesulfonate) (PSS). The modified GFs were then incorporated into a polypropylene matrix. It was also tested the hypothesis that the LbL technique would allow the incorporation of nanocomponents (such as carbon nanotubes) within the LbL layers. These nanocomponents are useful for introducing new functionalities on composites, such as to increase the stiffness, thermal and electrical properties[5,12,13]. Mechanical properties will be assessed in this work aiming structural applications.

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Layer-by-Layer technique employed to construct multitask interfaces in polymer composites

2. Materials and Methods

2.4 LbL modification of glass fibers

2.1 Materials

For the coating process, untreated glass fabric was used as the substrate for the LbL procedure. LbL steps are illustrated in Figure 1. First, glass fiber fabrics were immersed in the PDDA 2 wt% cationic solution for 5min.

Then, samples were rinsed with DI water for 2min and slightly dried under nitrogen stream for 1min. Reduced washing stages (even less than 2min) are commonly used on LbL procedures[10,17,18]. After the deposition of the first layer positively charged, the sample was immersed in the 0.30mg.ml-1 MWCNT-COOH/PSS anionic suspension for 5min, followed by a new wash and dry cycle. From the third deposition cycle on, the immersion time in cationic and anionic suspensions was reduced to 2min. Two distinct procedures were performed, with deposition cycles of 5 bilayers and 20 bilayers. The glass fibers resulting from those procedures will be referred as GF-5BL and GF-20BL.

Poly(diallyldimethylammonium chloride) (PDDA) solution 20 wt% in water was purchased from Sigma Aldrich. Average molecular weight of PDDA was 200,000-350,000 g mol-1.

The solution was diluted to 2 wt% using deionized (DI) water (18MΩ.cm). Poly(sodium 4-styrenesulfonate) (PSS) in powder was purchased from Sigma Aldrich.

Average molecular weight of PSS was 70,000 g mol-1.

DI water (18MΩ.cm) was used for preparing a solution of PSS 2 wt%. Multiwall carbon nanotubes carboxylic acid functionalized (MWCNT-COOH) were purchased from Sigma Aldrich and produced by Nanocyl (Belgium).

The nanotubes had an average diameter of 9.5nm, average length of 1.5 µm and carbon basis >80%. Functionalization was superior to 8 wt%. Bidirectional (0°/90°) glass fiber fabric was purchased from Casa da Resina, located in Belo Horizonte, Brazil. It was a woven glass fabric (E-type glass fibers having density of 2.60g/cm3) with a specific mass of 120 g/m2, produced by Owens Corning. It will be referred in the text as commercial fibers. Polypropylene homopolymer PH 0950 was purchased in pellets from Braskem. Density of the polypropylene was 0.90g/cm3 according to manufacturer.

2.2 Preparation of MWCNT-COOH suspension A quantity of 300mg of carbon nanotubes functionalized with carboxylic acid groups was dispersed in 1L of aqueous solution of PSS 2 wt%[14]. The suspension was sonicated for 2h and then placed for 30min on tip sonicator, 20kHz, 40% amplitude (180W) with intermittent cycle 40s on / 20s off.

Those parameters were previously evaluated and adjusted in order to get the lowest sonication time while maintaining the sample homogeneous. This was tested consecutively every 10min until achieving a homogeneous dispersion. Ice bath was used to avoid suspension heating. A 10ml fraction was placed on total rest for qualitative analysis of the suspension stability during 4 weeks.

2.5 Production of bidirectional composites Pure polypropylene (PP) sheets of 0.45mm in thickness were fabricated by compression molding at 220°C and 30bar using a CARVER 4386 press. Composites were prepared by compression molding using a steel mold with cavity dimensions of 80mm x 11mm x 2mm. Initially, five PP sheets were interlaid with four layers of glass fiber fabric on the cold mold. The press was preheated to 210°C. The mold was taken to the hot press and a minimum pressure (close to zero) was applied for 2min30s. The pressure was relieved to zero and returned to contact condition for 30s. After that, the pressure was increased to 5bar every minute up to 30bar, remaining for 1min30s. The mold was then removed from the hot press and cooled during 15min under cold press (SAGEC) to 25°C and 20bar.

2.6 Characterization The scanning electron microscopy (SEM) images of GFs were obtained on a Quanta FEG 3D FEI system at an acceleration voltage of 5kV. A 5nm platinum coating was sputtered onto the surface of GFs to minimize the charging

2.3 Pre-treatment of glass fibers The commercial glass fiber fabrics were cut to the dimensions 180mm x 240mm and heated up to 650°C for 1h. After that, samples were immersed in a nitric acid solution (HNO3, 20% v/v) at 25°C for 24h and dried at 60°C for 24h in order to remove the commercial sizing.

Subsequently, samples were immersed in 600ml of hydrogen peroxide solution (30% v/v) at 75°C for 45min.

30ml of ammonium hydroxide solution (30% v/v) was added dropwise. After 5min, samples were washed with DI water and dried on oven at 60°C for 2h. The purpose of this step was to increase the concentration of hydroxyl groups on the surface of the fibers[15]. The glass fibers from this procedure will be referred in the text as untreated glass fibers (GF-untreated). No characterization was performed for the commercial sizing since the main interest at this stage rested on its elimination to start the coating procedure[16].

Polímeros, 27(4), 330-338, 2017

Figure 1. Representation of the Layer-by-Layer procedure.

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Vitorino, L. S., & Oréfice, R. L. effects. The cross sections of modified GFs were obtained by transverse cutting. Individual glass fibers were mounted on cardboard frames and were cut off using a saw blade to get a plane perpendicular to the fiber direction. Moreover the modified GFs were assessed by focused ion beam scanning electron microscopy (FIB-SEM) for evaluation of the coating thickness. The fibers were FIB-milled with a gallium (Ga+) ion beam current of 1nA and with the acceleration voltage of 30kV. The ion beam was used for milling, not for imaging.

The cross-sectioned face was polished with the low beam current prior to imaging with SEM since Ga+ ions could be deposited in the cross sectioned wall impeding its perfect visualization.

Thermogravimetric analyses (TGA) for the untreated and modified glass fibers were performed on an instrument EXSTAR TG/DTA 7200 with air atmosphere, 30mL/min flow and a heat rate of 10°C/min.

Raman spectra for the MWCNT-COOH and the LbL modified glass fibers were obtained on a Jobin Yvon/Horiba LABRAM-HR 800 spectrograph equipped with a He-Ne laser (632.8nm). The Raman signal was collected by a microscope Olympus BX-41 provided with objectives (10x, 50x and 100x). The detector used was a N2 liquid cooled CCD of Spectrum One, back illuminated. Depending on the sample background fluorescence, the acquisition time ranged from 10s to 120s and the laser power from 0.08mW to 8mW. To improve signal/noise ratio, spectra were acquired 10-30 times.

Density measurements of the composites were done accordingly to ASTM D792 - Method A, using water as immersion liquid. The constituents content of the composites were obtained according to ASTM D3171 (procedure G, matrix burnoff in a muffle furnace) at 650°C during 1h.

Calorimetric analysis of the composites using differential scanning calorimetry (DSC) was performed by EXSTAR DSC 7020 equipment at the heating rate of 10°C/min. Tests were conducted on 9-10mg samples from 20°C to 220°C (first heating) and melting enthalpy and crystallization degree

were evaluated. Relative crystallinity (Xc) was calculated as the ratio between the value of the melting enthalpy and the theoretical value of the melting enthalpy for the completely crystalline polypropylene (207J/g)[7]. Polypropylene directly impregnated on the bidirectional glass fiber fabric containing few fiber segments was extracted from the composite plate in order to get the DSC samples.

Flexural measurements for the bidirectional composites were done according to ASTM D790 - Procedure A, flat wise direction. Data were collected from minimum 5 specimens which were cut and prepared from the composite plates.

The specimen dimensions were 12.7mm x 50.0mm x 2.0mm.

A universal testing machine INSTRON 5965 equipped with a 5KN load cell and crosshead speed of 0.90mm/min was used. The samples were cryo-fractured under liquid nitrogen (30min immersion) and the fracture surfaces of the composites were studied using a scanning electron microscope (FEI Model Inspect S50) at an acceleration voltage of 15kV. A 5nm platinum coating was sputtered onto the surface to minimize the charging effects.

3. Results and Discussion 3.1 MWCNT-COOH suspension stability Before sonication of the carbon nanotubes in PSS, aggregates were initially being deposited into the bottom of the beaker. After 2h30m of sonication, the suspension became totally dark and homogeneous. No aggregates were found on the bottom of the beaker after a closer inspection.

The appearance has remained the same after four weeks without visible aggregates or sedimentation. Figure 2 shows the dispersion state of the sample as a function of time.

Carbon nanotubes in suspension represent a kinetically stable system rather than thermodynamically. Thus, sedimentation and aggregation may occur over time[19]. Since the dispersion was an intermediate stage in this work and it was used immediately after in the LbL deposition step, the observed stability may be sufficient. The good dispersion of carbon

Figure 2. Dispersion state of MWCNT-COOH/PSS suspension as function of time.

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Layer-by-Layer technique employed to construct multitask interfaces in polymer composites nanotubes in PSS can be explained by the presence of aromatic rings which have a strong affinity for graphitic surfaces via π-stacking. Typical effective dispersants are polymers consisting of repetitive units of an alkyl chain and aromatic rings, exploiting both the π-stacking on the surface of CNTs and the Van Der Waals attractions between the hydrophobic part of the surface of nanotubes and the alkyl parts[20].

3.2 Glass fibers characterization The pre-treatment relative to sizing removal was evaluated

by SEM by comparing the morphology of commercial glass fiber with the untreated glass fiber. Figure 3 shows the images of commercial (Figure 3a) and untreated (Figure 3b) glass fibers. The average fiber diameter is 15μm. The presence of a homogeneous coating is observed over the entire surface on the commercial fiber. After the cleaning treatment, the surface appearance became smooth and clean. The straight lines seen on the borders indicate the absence of a coating.

The LbL procedure (5BL and 20BL) changed the cleaned glass fiber fabric coloration from its original white to dark gray, that could be associated with the deposition of the carbon nanotubes. The 20BL sample became distinctly darker, suggesting a larger amount of deposited carbon nanotubes.

Moreover, the water used during the washing step presented a greyish coloration during the first minute of washing, indicating that weakly adsorbed carbon nanotubes were removed and thus preventing cross-species contamination.

After the LbL procedure, 5BL and 20BL modified glass fibers were analyzed by SEM as also illustrated in Figure 3.

It is possible to observe the presence of carbon nanotubes on GF-5BL surface. The general appearance is still smooth, preserving the straight lines on the borders similar to the untreated fibers. This could be associated with the very thin coating seen in the sectioned view obtained by FIB milling (Figure 4a) with thickness 100-200nm. Meanwhile, the surface morphology has dramatically changed for 20 bilayers deposition. A high increase on surface roughness is

Figure 3. SEM images of glass fibers: (a) commercial; (b) untreated; (c), (d) GF-5BL; (e), (f) GF-20BL. Polímeros, 27(4), 330-338, 2017

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Figure 4. SEM-FIB images of cross sections of FIB milled fibers: (a) GF-5BL; (b) GF-20BL.

Figure 5. SEM images of cross sections of mechanically cut fibers: (a) GF-5BL; (b) GF-20BL.

observed with the formation of a coating along all the fiber. The cross view of the GF-20BL after FIB milling is shown in Figure 4b. The thickness of the coating could be assessed as 400-500nm. It was not possible to observe the individual carbon nanotubes since they were embedded in the thick polymeric coating.

Additional SEM images were acquired for glass fibers which were cross sectioned by transverse cutting (Figure 5).

It is possible to observe that the coating thickness initially assessed by localized FIB-milling on the glass fiber surface is actually variable along the fiber.

The nanocoating deposited by LbL method on glass fibers is a novel application. In this work the obtained values for the coating thickness have the same order of magnitude of values reported by other authors[5,8] although a different method is used.

Raman spectroscopy, a powerful tool for carbon nanotubes characterization, was employed to demonstrate the deposition of CNTs on GF surface. In Figure 6, characteristic peaks of CNTs are identified. The D band derived from defects in the nanotube wall is found at 1326cm-1. The G band which represents the crystalline graphitic and in-plane vibrations of sp2 carbon is displayed at 1580cm-1 and D* band which represents the overtone of disorder is found at 2645 cm-1[21].

The presence of the same Raman peaks on the modified GFs proved that carbon nanotubes were successfully incorporated 334

Figure 6. Raman spectra of GF-untreated, MWCNT-COOH and GF-20BL.

within the LbL layers. For the untreated glass fiber, no peak is observed.

Thermogravimetric analysis was employed to get information on the coating weight content. TG curves for the untreated and modified glass fibers are seen in Figure 7. For untreated and 5BL glass fibers, small mass Polímeros, 27(4), 330-338, 2017

Layer-by-Layer technique employed to construct multitask interfaces in polymer composites fluctuations are observed. For the untreated glass fibers, this TG result confirmed SEM results that showed that the removal of commercial sizing was successful via matrix burnoff and acid digestion. For 5BL glass fibers, a range of 100-200nm for the coating thickness was estimated by SEM but variations along and between fibers can occur, leading to a lower weight loss. Moreover, 100nm polymer based coatings in 15000nm thick glass fibers (with densities much higher than the density of the polymer coatings) result in a very low change of weight (around 1%) of the system that may be within the measurement error of the analytical technique (i.e. TGA).

TG curve for glass fibers coated with 20 BL showed 6.1% mass loss which can be assessed as an approximation of the mass content of LbL coating deposited on those fibers.

The value found is lower than the actual one since it is known that PSS burning leads to inorganic by-products of high stability due to the presence of Na+ in its structure[22].

3.3 Composites characterization 3.3.1 Physical and thermal properties The density values at 23°C, the GF content and the void volume obtained for the prepared composites are reported in Table 1.

Changes in density of a same material can occur due to localized crystallinity and/or to different proportions of resin

and glass fiber (ASTM D792). Comparing the density values among the composites, a mean of 1.072g/cm3 and a standard deviation of 0.001g/cm3 indicate good process repeatability on composites fabrication. Comparing the glass fiber content among the composites, a mean of 25.62 wt% and a standard deviation of only 0.39 wt% was found, corroborating the idea of good process repeatability. The influence of the fiber content on the flexural properties will not be taken into account since a small variation among the composites was obtained.

The obtained coatings on the fibers are constituted by the polymers PDDA and PSS and by the carbon nanotubes.

Densities of those constituents are similar to the density of the composite matrix (polypropylene, 0.9g/cm3). Considering also that the content of the coated fibers in the composites is relatively low (around 25 wt%) and that the coating of nanometric dimension is just at the fibers surface, the density of the composites will be not significantly influenced by the coating. The other parameters on Table 1 will not be influenced by the coating as well, since the procedure for obtaining fiber content and void volume is a burnoff in a furnace at 650°C which burns also the polymer content (including the polymeric coating).

The mean value around 0.9% for the void volumes indicates a good quality composite (ASTM D2734). The high standard deviation can be explained by the fact that as the void content gets lower the error in resin density gets increasingly more important. According to ASTM D2734, ideal measurements of density should have an uncertainty of less than 0.0005g/cm3, but a 0.005g/cm3 was found.

The relevant thermal properties (melting temperature peak TM,PEAK, melting enthalpy ∆HM and relative crystallinity XC) of the prepared composites are summarized in Table 2.

All composites presented a single endothermic peak with very similar positions, suggesting that the modification of the GFs did not affect the melting temperature of the composites. The similar values of the relative crystallinity (with 1% max. variation) show that the modification of the GFs did not influence the crystallization degree.

3.3.2 Morphology

Figure 7. TG curves for GF-untreated, GF-5BL and GF-20BL.

For glass fiber reinforced composites, the morphology of the fracture surface can give information related to the adhesion of fiber-matrix interface. High interfacial adhesion represents an essential key for obtaining high-performance composites[5-6].

Table 1. Density and constituents content for the prepared composites.

Composite PP-GF-Untreated PP-GF-5BL PP-GF-20BL

Density (g/cm3) 1.071 ± 0.004 1.071 ± 0.004 1.073 ± 0.006

GF content (wt%) 25.49 ± 0.13 26.05 ± 0.77 25.31 ± 0.40

Void volume (%) 0.83 ± 0.10 1.27 ± 0.60 0.51 ± 0.31

Table 2. Relevant thermal properties of the prepared composites.

Composite

Tm,peak (°C)

∆Hm (J/g)

Xc (%)

PP-GF-Untreated PP-GF-5BL PP-GF-20BL

165.8 165.9 166.2

64.0 66.0 65.0

30.9 31.9 31.4

Polímeros, 27(4), 330-338, 2017

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Figure 8. SEM images of the fracture surfaces of the composites: (a) PP-GF-untreated; (b) PP-GF-5BL; (c) PP-GF-20BL; (d) PP-GF-20BL 4,000x magnification.

The images obtained by SEM for all composites (Figure 8) indicate that the fracture occurred predominantly by interfacial debonding with many fibers pull-out. The fibers without coating (PP-GF-untreated) failed by progressive fiber pull-out and debonding. The extensively pulled-out fibers fractured with different lengths could indicate that the damage initiated from several locations along the interface[5].

The coated fibers failed transversely across the specimen with more uniform length distribution although debonding and pull-out are also the main fracture micromechanism.

The smooth and clean surface of the fibers without evidence of adhered resin suggests a low adhesion condition with the matrix and therefore a weak interface. The interface between polymer and GFs plays an important task in controlling some of the mechanical properties as tensile and flexural strenght.

The fiber-matrix adhesion is confined to the interface where stress-transfer shall occur[4-6]. Thus the results presumably indicate that the composites will not present an efficient load transfer from the matrix to the reinforcement phase (GFs) when subjected to mechanical stress.

For the glass fibers coated with 20 bilayers, the nanocomposite coating appears intact indicating a strong bond to the fiber and no changes due to the composite processing steps (Figure 8d).

3.3.3 Flexural mechanical properties Stress vs. strain curves for all the evaluated composites are shown in Figure 9. The curves were chosen from the specimen most representative of the mean results. Flexural Modulus and flexural strength are reported in Table 3. The flexural strength was calculated at 5% of strain since the samples did not show breaking before that strain.

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Figure 9. Flexural stress-strain curves for the prepared composites.

It can be noticed in Figure 9 that the presence of LbL layers on GFs modified the stress vs. strain curves. The total area under the curve assigned to the PP-GF-20BL composite is higher than the area under the other curves, meaning that the presence of the 20 bilayers was useful in increasing toughness of the material. This behavior can be attributed to a ‘healing’ effect in which GF’s micro cracks are filled by the coating and severe fiber surface flaws are eliminated due to an increased crack tip radius[5,23]. Moreover, the presence of the polymer coating at the interface can provide mechanism for energy dissipation during crack propagation.

Polímeros, 27(4), 330-338, 2017

Layer-by-Layer technique employed to construct multitask interfaces in polymer composites Table 3. Flexural mechanical properties for the prepared composites.

Composite PP-GF-Untreated PP-GF-5BL PP-GF-20BL

Flexural Modulus (MPa) 2730 ± 417 2921 ± 264 3174 ± 183

Flexural Strength (MPa) 56.4 ± 4.0 56.4 ± 2.5 57.5 ± 1.0

The mean values for the Flexural Modulus indicate an increase in this property as the number of bilayers rises (Table 3). The PP-GF-20BL composite exhibited an increase of 16.3% when compared to the composite with untreated fibers and 8.7% when compared to PP-GF-5BL. The elastic Modulus is measured at very low deformation and in this condition (close to zero strain) materials react separately to an atomic-molecular level. Thus, the Modulus prediction is not influenced by the low interfacial adhesion observed between fiber and matrix and the properties of each single component become more important[4]. We suggest that an improvement in the glass fibers individually due to the presence of carbon nanotubes on their surface substantially increased the stiffness of the system. Coatings containing carbon nanotubes on GFs are useful for enhancing fiber`s mechanical properties[5,8]. The coating`s Modulus and thickness can be also associated as reasons for the results[23].

Nevertheless it is known that mechanical damage on carbon nanotubes, i.e. breakage, induced by high-power tip sonication leads to defective carbon nanotubes with critical reduced length[24]. The preparation of the MWCNT-COOH suspension with 30min of tip sonication in this work was necessary to achieve a homogeneous and stable suspension in PSS. This process could impair the mechanical properties associated to the defective CNTs thus bringing the coating Modulus to smaller values. The Modulus values obtained for the final composites would also be influenced in a similar way.

There was no significant improvement on the flexural strength among the composites (Table 3). The mechanical strength is highly influenced by the interface adhesion since the maximum stress that material withstands relies on the capability of transferring load between the matrix and the fiber[4]. The results of SEM for the crio-fractured surfaces (Figure 8) indicated a low-adhesion condition between GFs and PP. Consequently it was expected that the flexural strength values for the modified GF composites became similar to the untreated GF composite due to ineffective load transfer at the interface.

Potential reason for the low adhesion is the lack of acid chemical groups on polypropylene capable to interact with electron donor sites on the fiber surface which was coated by the polar PDDA/PSS layers[21]. The LbL method makes use of polar constituents since it is based on cationic and anionic electrostatic interactions for the construction of the bilayers. On the other hand, the PP is a non-polar polymer.

Thus it is possible that this result was related to the lack of chemical affinity between the PP and the LbL components.

However, due to its versatility, the LbL technique can provide ways to overcome this observed lack of adhesion.

By changing the types of polymers deposited through the Polímeros, 27(4), 330-338, 2017

LbL technique, e.g. by using ionic polymers with alkyl groups in the last deposited bilayer[25,26], it could be possible to increase the affinity between the LbL coating and the polymer matrices in future works.

4. Conclusions In summary, carboxylic acid functionalized multiwall carbon nanotubes were dispersed in PSS aqueous solution and a homogeneous stable dispersion was achieved.

The dispersion was used as an anionic polyelectrolyte in LbL system to successfully deposit a nanocomposite coating (PDDA//MWCNT-COOH/PSS)n on glass fibers surface as proved by SEM and Raman characterization. TGA measurements have shown an estimated coating content on the fiber around 6 wt% for 20 bilayers. The nanocomposite coating has created a new complex interface between fibers and PP matrix when bidirectional composites were fabricated.

The flexural mechanical properties of the composites were analyzed, showing an increase of 16.3% in the Flexural Modulus and toughness for the PP-GF-20BL composite when compared to PP-GF untreated. Flexural strength was not affected by the GF coating. The results of this work proved that the LbL technique can be used to modify the interface of polymer composites and to incorporate nanocomponents (such as carbon nanotubes) within their interfacial region.

The possibility of tailoring the composition and structure of the interface in polymer composites through the use of the LbL technique can allow the construction of interfaces that can perform multiple tasks, such as increasing toughness and encapsulating active components.

5. Acknowledgements The authors would like to acknowledge CAPES, CNPq, FAPEMIG, INCT Acqua and Center of Microscopy of the Federal University of Minas Gerais for providing the equipment, technical and financial support for all experiments. Part of this research was conducted at FCA Fiat Chrysler Automobiles, Betim, Brazil.

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Received: Nov. 15, 2016 Revised: Mar. 09, 2017 Accepted: Mar. 14, 2017

Polímeros, 27(4), 330-338, 2017

http://dx.doi.org/10.1590/0104-1428.14916

Influence of tribological test on the global conversion of natural composites Carlos Eduardo Correa1, Robin Zuluaga2, Cristina Castro3, Santiago Betancourt4, Analía Vázquez5 and Piedad Gañán6* Grupo de Nuevos Materiales, Universidad Pontificia Bolivariana – UFB, Medellín, Antioquia, Colombia 2 Facultad de Ingeniería Agroindustrial, Universidad Pontificia Bolivariana – UFB, Medellín, Antioquia, Colombia 3 Facultad de Ingeniería Textil, Universidad Pontificia Bolivariana – UFB, Medellín, Antioquia, Colombia 4 Facultad de Ingeniería Mecánica, Universidad Pontificia Bolivariana – UFB, Medellín, Antioquia, Colombia 5 Consejo Nacional de Investigaciones Científicas y Técnicas – CONICET, Instituto de Tecnología en Polímeros y Nanotecnología, Universidad de Buenos Aires – UBA, Buenos Aires, Capital Federal, Argentina 6 Facultad de Ingeniería Química, Universidad Pontificia Bolivariana – UFB, Medellín, Antioquia, Colombia 1

*piedad.ganan@upb.edu.co

Abstract The vinyl ester resins and natural composites have emerged as a suitable alternative in tribological application due to mechanical behavior, which relates to the conversion of the double bonds. During tribological test the permanent contact between polymeric sample and counterpart can increase the temperature affecting the crosslinking of the samples. These variations have direct implications in the curing rate and the global conversion. In this work, the FTIR evaluation is used to evaluate possible changes on the global conversion of vinyl ester and their composites reinforced with Musaceae fiber bundles and cured using two hardeners, after a specific tribological test. Increments around 15% on global conversion of styrene double bonds were observed for neat matrix and composites using both hardeners, suggesting that during tribology test some alterations on resin structure takes place. These results open alternatives to manipulate the curing conditions in order to control the tribological behavior.

Keywords: FTIR analysis, global conversion, natural fiber composites, tribology test, vinyl ester matrix.

1. Introduction Polymers such as unsaturated polyester resins, epoxy resins and vinyl ester resins have been used as matrix for composite materials in an uncountable amount of technical applications[1-11]. The most popular of them due to the low cost, reasonable good properties and simplicity for processing are the unsaturated polyester resins, but when better mechanical or chemical resistance properties are required, the epoxy resins are preferred. However, they are commonly used in applications where the cost is not essentially a problem.

The vinyl ester resin, being a mixture of the previously mentioned resins combines the easy processing and curing of the unsaturated polyester resins with the high mechanical, thermal and chemical properties of the epoxy resins[9,11-13].

The vinyl ester resins are a kind of unsaturated polyester resins that have chain terminations of an acrylate or methacrylate groups. They are family with the common unsaturated polyester resins, but they usually are diesters that contain ether linkages and vinyl functional groups[8,14].

The main structure may come from several other common resins such as epoxy, urethane or many more, but the most common are the epoxy resins.

The resins with an epoxy backbone are commonly a copolymer of ethylenically unsaturated carboxylic acids

Polímeros, 27(4), 339-345, 2017

like acrilyc or methacrilyc acid[8,15-18] and components of the epoxy resins, for example the diglycidyl eter of bisphenol A (DGEBA). These resins may be used neat[15,19,20], however, they are usually diluted in a vinyl-type reactive co-monomer, most commonly used is the styrene monomer[8,14,15,19,21]. The reaction may be catalyzed by ammonium salts, ternary amines, alkalis or phosphines and it takes place at moderately elevated temperatures[8,10], temperatures between 60 to 120 °C. Curing of vinyl ester resins is potential by C=C double bond located at the ends of the chains. Due to this fewer C=C bonds these resins have a lower crosslinking density compared with the unsaturated polyester resins, and as these bonds are in the extremes of the chain, they are easier to be accessed, which may lead to an almost complete conversion in the curing[8,22].

The vinyl ester resins can be possibly used in less common applications as tribological evaluation, when it is always recommended that matrix experiment the less modification possible as weight loss. The potential of vinyl ester resins and their composite use in tribology are increasing, some works in this field include composites reinforced with

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O O O O O O O O O O O O O O O O

Correa, C. E., Zuluaga, R., Castro, C., Betancourt, S., Vázquez, A., & Gañán, P. glass reinforcements[23,24], SiC filler[23], carbon fabric[24], or natural fibers[25,26].

During the tribological test increments on temperature due to the contact with counterpart, traditionally metallic, are possible. These phenomena could affect the crosskling of vinyl ester samples. However, there are no studied about the effect of this test on the resin structure during or after this takes place, in spite the importance of the plastic behavior of the tribological systems.

This potential affection is based on that the curing process is a heterogeneous, highly exothermic free radical copolymerization[15,22] where three reactions occur at the same time: vinyl ester homo-polymerization, styrene (or any other unsaturated vinyl monomer as mentioned before) homo-polymerization and vinyl ester–styrene copolymerization[3,8,13,17]. The crosslinking reaction undergoes in differential form because the vinyl ester double bonds at the beginning of the reaction, reacts faster than the ones of the styrene, but in the final part is the opposite.

This differential in the crosslinking forms microgels which grows and coagulate and then adding between them, which brings to the final nodular form[5,13,22,27]. The reaction rate of the vinyl ester and the monomer may be divided in two parts, a faster one, where the maximum rate of reaction is located may be due to the easier diffusion of the components.

However at higher conversions a three dimensional gel structure appears and the diffusion of the long chain radicals is more difficult to occur and the reaction slows down[6], and potential alterations on temperature of the sample could affect this last conversion.

In order for curing process to take place, peroxides or hydroperoxides are used as initiators, and are also known as hardeners. At high temperature curing these hardeners may be used alone or at room temperature curing is required the use of a redox initiation system[3] of promoter (initiator, hardener) and activator (accelerator), these are systems of the peroxide, cobalt salts and/or ternary amines[28], not always with the three components. The vinyl ester resins may also be cured via photo-polymerization[8] using an appropriate photoinitiator[21]. Radiations used to the photo-curing of the resins are ultraviolet[14,29], visible light[7] and electron beam[14,30].

For the curing of the vinyl ester the hardeners used are methyl ethyl ketone peroxide (Mek peroxide) which is the most common for room temperature curing[6,20,29,31]. Also are used diacylperoxides, peresteres, diaryl peroxides, dialkyl peroxides, hydroperoxydes such as cumene hydroperoxyde[3] cumyl hydroperoxyde[5], for high temperature curing the most common initiator is benzoyl peroxyde (BPO)[1,12,15,19], but t-butyl proxy-benzoate[22] is used as well sometimes.

An activator does the decomposition in free radicals of the initiators. This activation may be done using chemicals, temperature or radiation. On the other hand, as secondary promoters a ternary amine is used to complement the action of the accelerator, or just used to accelerate when the hardener is BPO. When the resin cures, there are several things that may affect the curing of the resin such as the temperature and time, the diluent reactive monomer, the amount of hardener, accelerator and inhibition products.

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When the styrene content is changed in part or completely with α-methyl styrene the viscosity, gel time and peak curing temperature have the same behavior as the one mentioned with only styrene when the fraction of the α-methyl styrene is increased in the formulation of the resin[1,19]. The variation of the components of the initiation system, cobalt salt and peroxide, also affects the crosslinking reaction kinetics.

When the isothermal curing temperature is increased the reaction rates and overall conversion of the resin increases[32].

However the final conversion of the resin has a maximum at intermediate temperatures where it reaches almost complete curing, but at higher temperatures than those mentioned the final conversion does not reach that value[6].

The global or final conversion degree of the cure after a post-curing process in the vinyl ester resin is independent of the isothermal curing temperature in the material[33].

The postcure of the resins may also be done via radiation like microwaves[11]. However, no mention or studies of variations of global or final conversion of the resin after a test, such as tribological test, have been found.

On the other hand, the curing and poscuring evolution of vinyl ester resins have been studied using different types of techniques such as thermal analysis DSC test, DMA test, FTIR spectroscopy or RAMAN spectroscopy[34].

FTIR spectroscopy and its variations such Attenuated total reflection (ATR) could be an interesting technique of analysis to obtain the curing process and to obtain the initial information about the potential alteration of resin structure due to modifications or alterations of conversion during applications or test.

As mention before, the vinyl ester is a very versatile system, and it is possible to manipulate is global conversion by variation on temperature such the case of the tribological applications. In this study, the potential use of ATR-FTIR spectroscopy to monitor the potential variations of conversion of vinyl ester during a tribological test is analyzed, and interesting variations on global conversion of matrix have been observed. Additionally to neat matrix, natural composites reinforced with natural fibers such as Musaceae fiber bundles are considered. The use of this type of fiber bundles is due to the interesting qualities of their composites for the tribological applications, that have been evaluated in previous work[25,35].

2. Materials and Methods 2.1 Materials In this analysis a vinyl ester resin, reference Swancor 901-3 kindly supplied by Andercol S.A. was used. Two hardeners, methyl-ethyl-ketone peroxide (Mek) and benzoyl peroxide (BPO) were used.

The fiber bundles used in this work were extracted from the rachis of Colombian Musaceae plants, and were kindly supplied by Banacol S.A. Fiber length used in this work corresponds to 287 μm.

2.2 Composite fabrication Composite samples were prepared using a BMC fabrication technique. Resin, hardener and fiber bundles were mechanically mixed, then the mix was put in a mold and compressed at 1000 psi. The curing condition corresponds to Polímeros, 27(4), 339-345, 2017

Influence of tribological test on the global conversion of natural composites 100 °C during 1 h, these conditions were defined according with previous work[35]. This condition was chosen in order to explore potential variations of curing conversion of vinyl ester resin after the tribological test. All composites samples contained a 10 wt % of fiber bundles, whereas as the amount of hardener corresponds to 1.5 wt %. Figure 1 summarizes the fabrication process. The void content of neat and composites samples were less than 1%. Additionally, non-debonding problems were registered in samples after fabrication process.

2.3 Tribological test Five samples with 9 mm in diameter were cut and attached to a 10 mm long metallic pin with cyanoacrylate contact adhesive. Pins were machined to reach a diameter of 6.3 mm. The tribological test was performed using a pin on disc machine as the one schematized elsewhere[35].

According to this previous study, the most adequate parameters of the experiment test corresponds[35]: 200 m min−1 of speed, recorded distance corresponds to 3 km, and normal load to 4.9 N. The counter body used was a 1040 steel, a common material for the evaluation of the tribological behavior of polymers. Figure 2 shows a pin on disc machine.

Figure 1. Scheme of vinyl ester resin and composites fabrication and evaluation.

2.4 Infrarred spectroscopy Attenuated total reflection Fourier transform infrared spectroscopy (ATR–FTIR) Nicolet 6700 supplied by Thermo Scientific was used for FTIR analysis. The FTIR spectra were recorded on a Nicolet 6700 spectrophotometer in the 4000-400 cm−1 range using ATR. The spectra were recorded with a resolution of 4 cm−1 and an accumulation of 64 scans.

Evaluations were developed on cured samples before and after the tribological test.

To evaluate the final or global conversion of vinyl ester double bonds (ɑVE) and global conversion of styrene double bonds (ɑS) before and after tribological test, in order to evaluate potential alterations on curing, Equations 1 and 2 are used.

These equations are adapted from Ziaee and Palmese work[5]. aVE 1 – ( ABS F 945 cm−1 / ABS0 945 cm−1 ) × ( ABS0 830 cm−1 / ABS F 830 cm−1 ) (1)

Figure 2. Pin on disc machine.

Table 1. Convention of analyzed samples.

Type of material Neat matrix before tribological test Neat matrix after tribological test Neat matrix before tribological test Neat matrix after tribological test Composite before tribological test Composite after tribological test Composite before tribological test Composite after tribological test

Curing agent Mek peroxide Mek peroxide BPO BPO Mek peroxide Mek peroxide BPO BPO

Abbreviation

NMBT-Mek NMAT-Mek NMBT-BPO NMAT-BPO CBT-Mek CAT-Mek CBT-BPO CAT-BPO

aS 1 – ( ABS F 910 cm−1 / ABS0 910 cm−1 ) × ( ABS0 710 cm−1 / ABS F 710 cm−1 ) (2)

2.5 ABS0 represents the peak height of uncured matrix ABSF represents the peak height at end process, which corresponds after curing process or after tribologial test.

Vibration at 945 cm-1 corresponds to out of plane bending of C-H double bond of vinyl groups, and 910 cm-1 corresponds to double bonds of styrene monomer. To evaluate the global conversion is required vibration not affected by curing process. In this case corresponds to aromatic groups of vinyl ester located at 830 cm-1 and 700 cm-1 for styrene monomer. Table 1 summarizes the sample convention analyzed in this work.

Polímeros, 27(4), 339-345, 2017

3. Results and Discussions As Figure 3 shows, uncured vinyl ester resin spectra, characteristic vibration of vinyl ester resin corresponds to ester linkage due to formed carbonyl group through the transformation of epoxy group are observed at 1715 cm-1[36].

A broad vibration between 3628 and 3142 cm-1 is associated with hydroxyl groups; centered peak corresponds to 3445 cm-1.

In areas around 1630, 945 or 830 cm-1 present vibrations correspond to C=C vibration of methacylate and styrene monomer, the main groups responsible of curing process of this resin. These vibrations can be overlapped. Characteristic vibration of aromatic ring are observed at 1607 and 1508 cm-1[34].

341

Correa, C. E., Zuluaga, R., Castro, C., Betancourt, S., Vázquez, A., & Gañán, P. Table 2 summarizes the main vibrations of the uncured vinyl ester resin dissolved in styrene.

FTIR technique helps to evaluate global conversion of monomers or reactive groups that participate during curing reaction, such as the case of styrene and vinyl ester double bonds[5]. In this case, global conversion before and after tribological test of these groups are evaluated for both neat and Musaceae fiber bundles composites. These composites were evaluated due to the potential importance for tribological

Figure 3. Uncured vinyl ester resin.

applications[25]. Figures 4 to 6 show the final conversion of final or global conversion of vinyl ester double bonds and global conversion of styrene double bonds before and after tribological test when two different types of hardeners are used. According with Figure 4, increments on global conversion of vinyl ester groups of neat resin samples are observed after tribological test. These modifications are most significant when the Mek-peroxide is used as hardener.

For styrene groups, the variations on global conversion of neat resin samples are registered in Figure 5. In this case, when BPO peroxide is used, the alterations are of the most significant. When composite samples are analyzed, see Figure 6, the global conversion of styrene double bonds groups before the tribological test is lower than neat resin, because composite samples presente at increments on the global conversion respect of the neat matrix. This behavior suggests that the presence of fiber bundles could affect the development of curing. Some authors[39,40] have commented alteration on curing evolution of thermosetting composites systems, even when natural fibers are used as reinforcement.

Musaceae fiber bundles could introduces the variations on the interfacial interactions respect to the matrix and introduce a catalytic effect on curing process and polymer chain network formation[41,42]. This phenomenon could introduce topological restrictions that reduce the impact of the evolution of the global convertion of the double bonds still present on composite samples after tribological test. Respect to the influence of tribological test on global conversion, a comparable behavior of neat resin (see Figure 5) is observed.

Table 2. Main vibrations of uncured vinyl ester resin.

Vibration (cm-1) 3626-3152 3079 3005 2924 2870 2851 1715 1636 1633 1609 1607 1581 1508 1297 1264 1244 1180 1041 1012 945 907 830 696

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Structure

References

OH groups of methacrylate groups. Broad region C-H aromatic ring C–H stretching vibration of alkene –CH2 and –CH3 groups –CH2 and –CH3 groups –CH2 and –CH3 groups Ester vibration Shoulder associated with C=C bonds styrene C=C bonds of methacrylate group Acryl double bond (–C=C–) Stretching of C=C present in aromatic ring C=C stretching of the aromatic ring Stretching of C=C present in aromatic ring C-O and C-C stretching Asymmetric vibrations of the C–O–C linkage and to the C–OH vibration Related with aromatic ring and C-O interaction Aromatic ring in plane and C-O interaction Aromatic ring in plane and C-O interaction Aromatic ring in plane Out-of-plane bending of vinyl ester monomer.

[33] [27] [9] [27] [27] [27] [36] [30] [30] [36] [37] [27] [37] [37] [36] [37] [37] [37] [37] [9,37]

C=C vibration of methacrylate C=C vibration presents in aromatic ring.

C=C in styrene Out of the plane vibration of C=C in aromatic group Out-of-plane bending of the CH groups in the aromatic ring, all the 5 hydrogen carbons oscillating in phase. Typical for monosubstituited aromatic ring

[30] [36,37] [38]

Polímeros, 27(4), 339-345, 2017

Influence of tribological test on the global conversion of natural composites

Figure 4. Global conversion of vinyl ester bonds of neat matrix samples before and after tribological test.

Figure 5. Global conversion of styrene double bonds of neat matrix samples before and after tribological test.

After the tribological test increments on global conversion are registered, essentially when BPO peroxide is used.

The main alterations are observed in S double bonds than in VE double bonds. This behavior has been observed by authors as Brill and Palmese[13], and potential explanation for that, is apparently that the VE reaction starts before the styrene one. However, the reaction of styrene continues after the VE one concludes. Additionally, these authors mentioned the formation of microgel or nanogel of VE that suggest the presence of multiple phases at the molecular level during the cure process. The results obtained in this work suggest that the tribological test affects this molecular distribution and potentially affects the mechanical or physical behavior[5].

The formation of these microgels are documented by other author such as Dua et al.[43] and Rodriguez et al.[27] when studied the curing of these resins at different isothermal curing. However, the results presented in this work are the first to document alterations of curing after a tribological test.

These results suggest that is necessary evaluate the curing alteration that occur in vinyl ester resins when are used for this type of applications, because the slight alteration could be manipulated to improve their behavior. The variations observed when samples are cured using BPO-peroxide respect to Mek-peroxide have relation with crosslinking evolution, and as well commented for other authors when analyze mechanical or thermal behavior[44].

These results are relevant considering that during tribological test potential plastification of the samples could affect the behavior of samples. If it is possible to control the evolution of the conversion of samples, important variations on tribological behavior could be obtained and unlimited potential applications could be developed even for neat matrix as well their natural composites because the alterations on the global convertion affects the plastic deformation of the systems, and it is reflected on the formation of film on the counterface of the counter part alterating the friction behavior.

These results could be useful for sliding applications or solid lubrications requirements at dry friction conditions.

4. Conclusions

Figure 6. Global conversion of styrene double bonds of natural composite samples before and after tribological test.

Polímeros, 27(4), 339-345, 2017

In this work, variations on the global curing conversion of vinyl ester resin and their natural composites before and after tribological test are analyzed using ATR-FTIR spectroscopy. According with obtained results, increments around 15% on global conversion of styrene double bonds are easier to follow using this technique. Increments superior to 30% were observed when BPO-peroxide is used as hardener, whereas the global conversion of vinyl ester double bond groups experiment more alteration when Mek-peroxide is used. The results suggest that ATR-FTIR spectroscopy could be an easier technique to use in order to monitor changes during important type of test such as tribological. However, the most important results of this work represent the relationship that exists between global conversion of double bond groups present in vinyl ester sample and tribological evaluations, according with it, is possible to manipulate the curing conditions in order to obtain different resin or matrix structure and the respective behavior during this mechanical test. These results are 343

Correa, C. E., Zuluaga, R., Castro, C., Betancourt, S., Vázquez, A., & Gañán, P. relevant considering that during tribological test a potential plastification of the samples could affect the behavior of samples. The alterations on the global convertion affects the plastic deformation of the systems and surface wear mechanisms. It can be reflected on the formation of film and transfer layer attached into the surface of the counterbody, both modifying friction behavior of tribosystem. In addition, if it is possible to control the evolution of the conversion of samples, important variations on tribological behavior could be obtained and unlimited potential applications could be developed even for neat matrix as well their natural composites, for instance, mechanical components in sliding conditions as gears or sliding dry pads in braking systems.

5. Acknowledgements The authors would like to thank CIDI and Colciencias for support and funding to develop this work. Also, thanks to Andercol S.A., and Banacol for providing materials used in this work.

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Received: Dec. 21, 2016 Revised: Mar. 05, 2017 Accepted: Mar. 21, 2017

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http://dx.doi.org/10.1590/0104-1428.12916

O O O O O O O O O O O O O O O O

Effect of concentrations of plasticizers on the sol-gel properties developed from alkoxides precursors Sandra Raquel Kunst1, Marielen Longhi1, Lilian Vanessa Rossa Beltrami2*, Lucas Pandolphi Zini1, Rosiana Boniatti2, Henrique Ribeiro Piaggio Cardoso2, Maria Rita Ortega Vega2 and Célia de Fraga Malfatti2 Centro de Ciências Exatas e Tecnologia – CCET, Universidade de Caxias do Sul – UCS, Caxias do Sul, RS, Brazil 2 Laboratorio de Pesquisa em Corrosão – LAPEC, Departamento de Metalurgia – DEMET, Universidade Federal do Rio Grande do Sul – UFRGS, Porto Alegre, RS, Brazil

1

*lvrossa@yahoo.com.br

Abstract Coatings developed through sol-gel method is presented as an interesting replacement to chromium coating. Sol-gel method present advantages as high purity and excellent distribution of the components. The objective of this work is to synthesize and characterize a film obtained by sol-gel route. The film was prepared with 3-(trimethoxysilylpropyl) methacrylate (TMSPMA), tetraethoxysilane (TEOS) and cerium nitrate, using water and ethanol as solvents. Polyethyleneglycol (PEG) plasticizer was added at four different concentrations. The sol was characterized by techniques of viscosity, thermogravimetric analysis (TGA), X-ray diffraction (XRD) nuclear magnetic resonance spectroscopy (NMR) and Fourier transform infrared spectroscopy (FT-IR). The results showed that tetrafunctional alkoxides condensation was retarded by the plasticizer, forming a compact film. The film with 20 g.L-1 of PEG showed the best electrochemical behavior.

Keywords: sol-gel, hybrid film, TEOS, TMSPMA, PEG.

1. Introduction Coatings prepared through sol-gel method is an ecological innovative technology due to anticorrosive properties of surface protection of metallic substrates, besides interesting properties in mechanic area[1].

Surface treatments still used in the packing industry are based on chromates, as they provide excellent corrosion resistance. Nontoxic alternatives to pre-treatments was developed in recent years to replace the chromating process.

Hybrid films obtained by sol-gel method are a good alternative to suppress the chromate-based process[2]. However, the industrialization of the sol-gel process requires testing with the real work substrates[3].

The sol-gel method promotes the polymerization of the precursors alkoxides silicon to forming a crosslinked network in hybrid films. The sol-gel method polymerization reaction is a two-step process. The first step is the hydrolysis of the alkoxy groups; the second step is the silanol condensation to form siloxane bonds.

Furthermore, the sol-gel obtained coating may have a heterogeneous roughness. In this case, it is recommended to coat with a thicker layer in order to improve the barrier effect or to coat a homogeneous and uniform tinplate substrate.

That thickness increase may be done in two ways: increasing the number of layers within limits to avoid delamination[4] or increasing sol viscosity. Sol viscosity can be controlled by temperature variation or by plasticizer addition. Temperature alters the hydrolysis and condensation reactions kinetics[5].

Hybrid films obtained by the sol-gel method may not be effective barriers against corrosion because they may present defects. To increase barrier protection, plasticizers, such as polymethylmethacrylate (PMMA)[6] or polyethyleneglycol

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(PEG)[7], have been used to increase coating thickness and to improve the flexibility of the system, which promote the acquisition of uncracked films that can be subjected to mechanical deformation without failure[8].

Generally, siloxane hybrid films are homogeneous and present good chemical and thermal stability. Additionally, these coatings have excellent barrier properties that improve the corrosion protection of pretreated substrates; these films act primarily as a barrier-type layer between the substrate and the environment. Consequently, they reduce the entry rate of water, electrolytes and oxygen and inhibit the permeation of species to the interface of the metal, which decreases the corrosion rate of the substrate. Accordingly, the degree of hydrophobicity and adhesion to the substrate (formation of strong covalent bonds; MeOSi)[6,7] are important properties of these films.

The technology of the sol-gel process is widely used to obtain high quality surface protection through simple procedures and with economic viability. Among its advantages, it is possible to mention: (I) the stoichiometry is easy to control and adjust[9]. (II) The production of a high purity film and with uniform distribution of its components[10]. (III) The process can be carried out under normal pressure and low temperatures.

Along the last decades a large number of hybrid materials has been obtained from the sol-gel process using various polymers and inorganic precursors[11-13].

In a preliminary study conducted by Kunst et al.[14], the tinplate was coated with a hybrid film composed of TEOS, TMSPMA, cerium nitrate in a concentration of 0.01 M and PEG 1500 in a concentration of 20 g.L–1. The authors tested the aplication of the monolayered and bilayered hybrid films

Polímeros, 27(4), 346-352, 2017

Effect of concentrations of plasticizers on the sol-gel properties developed from alkoxides precursors at different cure temperatures. The results showed that the bilayered hybrid film obtained at 60 °C had a higher layer thickness, and the best performance in the electrochemical assays, as well as the most hydrophobic character, in relation to the other samples. For the monolayered systems, the 90°C-cured system showed a lower layer thickness; however, this system showed a more compact, uniform and less porous layer, and presented better electrochemical impedance results, in comparison with the 60 °C-cured samples[14].

The objective of this work is to prepare a hybrid film obtained from sol constituted by 3 - (trimethoxysilylpropyl) methacrylate (TMSPMA) and tetraethoxysilane (TEOS) and with different polyethylene glycol plasticizer concentration (PEG). This paper evaluated the influence of PEG concentration on the sol formulation in the final physical and chemical properties of the films.

2. Experimental 2.1 Synthesis of hybrid films The hydrolysis reactions were conducted with the silane precursors 3-(trimethoxysilylpropyl) methacrylate (C10H20SiO5, TMSPMA) and tetraethoxysilane (C8H20SiO4, TEOS), with 0.01 mol.L-1 cerium nitrate (Ce III) addition (approximately 0.5% by weight) as corrosion inhibitor.

Ethanol and water were used as solvents.

The hybrid sols used as the coatings were prepared with the following molar ratios: TEOS: TMSPMA = 2:1; H2O:Si = 3.5:1; ethanol:H2O= 1:2. Polyethyleneglycol (PEG, 1500 g.mol-1) was added to the sol formulation at four different concentrations: 0, 20, 40, 60 and 80 g.L-1 (TP-0, TP-20, TP-40, TP-60 and TP-80, respectively) and TR is tinplate without coating. Ethanol and water were used as solvents. The hydrolysis time was 24 hours. According to further characterization tests, the hybrid films were prepared on a petri dish or applied by dip-coating, with 10 cm.min-1 removal rate and 5 minutes of immersion. Posteriorly, the samples was thermally cured at 60 °C for 20 minutes in a furnace. Figure 1 presents the flowchart of the process.

2.2 Experimental techniques Sol viscosity was measured after 24, 48, 72 and 96 hours of hydrolysis by a Brookfield DV2T viscometer with variable rotation and constant temperature (25 °C) for all samples.

The TGA were performed with a 50-Shimadzu TGA instrument over a temperature range from 23 to 700 °C and with a heating rate of 10 °C.min-1 under a nitrogen (N2) flow rate of 50 mL.min-1.

Free standing films were studied by XRD and FTIR. X-ray diffraction (XRD) analysis was performed on a Shimadzu XRD-6000 X-ray diffractometer, using a 2θ setting between 10° and 55°, with Cu Kα radiation of 1.5406 Å and step of 0.05°. The analysis of Fourier transform infrared spectroscopy (FTIR) was carried out using the attenuated total reflectance (ATR) on Nicolet IS10 Termo Scientific equipment. Each spectrum was obtained by performing 32 scans between 4000 cm–1 and 400 cm–1.

Solid Nuclear Magnetic Resonance (NMR) analysis of the hybrid films was carried out in an Agilent DD2 500/54 with a magnetic field of 11.7 T (500 MHz for 1H). 13C and 29Si experiments were conducted. A rotation of 10000 Hz, a pulse of 2.55 µs, a delay of 5 s and a contact time of 7 ms were the parameters for the 13C experiment. For 29Si, a rotation of 5000 Hz, a pulse of 3.2 µs, a delay of 5 s and a contact time of 9 ms were set. Furthermore, a liquid 13C NMR analysis was performed for a PEG 1500 sample in deuterated chloroform, in an Agilent 400 MR with a magnetic field of 9.4 T (400 MHz for 1H) and using TMS as internal standard.

The infrared spectroscopy measurements were performed using the technique of attenuated total reflectance (ATR), on Nicolet IS10 Termo Scientific equipment. The measurements were performed with the mid-infrared and each spectrum was obtained by performing 32 scans between 4000 cm–1 and 400 cm–1. The spectra were obtained for the films without a substrate (free-standing films).

3. Results and Discussion 3.1 Sol viscosity The fluid viscosity is regarded as the spontaneous creep or even to flow resistance and is due to internal friction (forces of attraction between molecules). Verification of viscosity is important to assess the integrity of the sol, being a valuable tool in quality control in the industrial field.

Figure 2 shows the variation of viscosity as a function of hydrolysis time samples. An increasing trend of viscosity as a function of the PEG concentration in the sol is observed, in which the TP-80 sample showed the highest viscosity.

Figure 1. Scheme illustrating the stages of the hybrid films development.

Polímeros, 27(4), 346-352, 2017

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Kunst, S. R., Longhi, M., Beltrami, L. V. R., Zini, L. P., Boniatti, R., Cardoso, H. R. P., Vega, M. R. O., & Malfatti, C. F. This behavior indicates that the higher the concentration of PEG in the sample, the higher the viscosity.

In addition, the increase of viscosity with time was linear.

The sol is a colloidal particles dispersion, that are stable in the fluid, and with time, they become gel formed by the rigid structure of colloidal particles or polymeric chains that immobilizes the liquid phase inside its interstitials[15].

This transition at room temperature, transforms the sol in gel by setting chemical bonds and molecular interactions between the particles. It allows the formation of a tridimensional solid network that provokes the increase of viscosity in the system up to the gelation point.

As shown in Figure 3, a solution containing water and PEG increases the viscosity gradually with the increase of the PEG concentrations. Low molecular weight PEG is more soluble in water and in nonpolar solvents. Moreover, most silanes hydrolyzable groups have limited solubility in water.

Until, though, these groups are converted to hydrophilic silanol groups by hydrolysis. The degree of polymerization of the silanes is determined by the amount of water available and the organic substituent. If the silane is added to water and has low solubility, a degree of polymerization is disfavored and subsequently has a lower viscosity as compared to the polyethylene glycol in aqueous solution[16,17].

Figure 2. Viscosity at 25 °C for the sols studied with different plasticizer concentrations.

3.2 Thermal characterization Figure 4 shows the thermograms of TGA of the studied hybrid films and the PEG. For the PEG in Figure 4 is observed only one stage of degradation, with a gradual and uniform weight loss. This event starts from 400 °C and ends at 422 °C. The PEG thermal degradation occurs by thermal cracking of the polymer chain.

The hybrid film without PEG (TP-0) presented a continuous process of weight loss, with a lot of waste at the end of the test. It was observed a loss weight at 110 °C, it was relative at loss of water residual and molecules absorbed in the film.

Subsequently, there has been a continuous weight loss up to the end of the test. This event relates to thermal degradation of TEOS and coupled to degradation of TMSPMA, both thermal main chain scission[18].

Figure 3. Viscosity at 25 °C, in water after 24 hours, for the solutions containing different PEG concentrations.

For hybrid films with PEG, an influence of the plasticizer on the thermal stability is observed. The addition of PEG improved the thermal stability of the silane coating when compared to the sample without PEG (TP-0). The samples with PEG showed a similar thermal behavior. These results demonstrate that the addition of PEG increases the thermal stability of the hybrid films and that the increase of PEG concentration favors this property.

3.3 Physic-chemical characterization Figure 5 shows the XRD diffractograms of the studied hybrid films. The silanes do not show arrangements in their crystalline structure, only a broad peak characteristic of the amorphous phase was observed. A gradually increase in the concentration of PEG in the samples is observed in Figure 5, it identified by increase in peak intensity. However, it was not possible to evaluate a change in the crystalline fraction of the hybrid films, since it is not possible to quantitatively compare the change in the intensity of the XRD peaks.

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Figure 4. TGA curves of the studied samples and PEG. Polímeros, 27(4), 346-352, 2017

Effect of concentrations of plasticizers on the sol-gel properties developed from alkoxides precursors Similar results were observed by Yang et al.[19] in their study on the effect of the concentration of PEG 1500 in composites with SiO2. The XRD showed that PEG 1500 has the crystalline structure with typical diffraction peaks at 19° and 24° (2θ). When PEG is added in greater concentration in the composite, it promotes a shift in the peak of the amorphous composite, increasing gradually intensity with increasing PEG concentration.

NMR analysis supplies information about the hybrid structures. It was carried out using MestreNova software.

Figure 6 shows the chemical structures of the silane precursors and of the plasticizer. 13C and 29Si spectra of the hybrid films appear in Figures 7 and 8. According

to 13C spectra (Figure 9), there was the formation of ether (C-O-C) function (4 in Figure 6) as the peaks at 66.57 ppm for TP-0 and 69.49 ppm for TP-20, TP-60 and TP-80 indicate (Solomons, 2011). This peak increases its area due to plasticizer addition, since PEG structure corresponds to that of ether, according to 13C NMR liquid PEG spectrum (Figure 9).

It is apparent the presence of ester (COOC) function, whose peak is located between 165 and 172 ppm[20]. This indicates that not all the TMSPMA participated in the hydrolysis process of this group. In addition, this peak area increases with the increase of PEG concentration, which allows us to infer that PEG enables molecular interactions, in this case, hydrogen bonds regarding the oxygen atom presence.

Peaks between 119-140 ppm reveal that the functional group C=CH2[20] relates to unpolymerized acrylate structures[21].

This peak reduces its area with the plasticizer addition.

The different trifunctional (Ti) and tetrafunctional (Qi) silicon-based structures appear in Figure 10[21-23]. Table 1 shows the peak localization in ppm for 29Si NMR spectrum of the different Ti and Qi species present in the hybrid structures.

Table 2 shows the polymerization degrees obtained of the hybrid film samples.

Each Ti and Qi proportion was calculated by integration of the NMR peaks, using MestreNova software. The relative proportions were calculated as follows (Equations 1 and 2)[22]:

Figure 5. XRD diffractogram of the hybrid films with different PEG concentration.

T i relative proportion =

Ti ∑T i

×100 (1)

Figure 6. Chemical structures of the silane precursors and the plasticizer.

Figure 7. NMR 13C spectrum for the hybrid films.

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Kunst, S. R., Longhi, M., Beltrami, L. V. R., Zini, L. P., Boniatti, R., Cardoso, H. R. P., Vega, M. R. O., & Malfatti, C. F.

Figure 8. NMR 29Si spectrum for the hybrid films.

Table 1. NMR 29Si peak localization of each silicon-based species in the hybrid films.

Hybrid Film

T1 -50.21 -48.97 -46.76 -57.94 -

TP TP-20 TP-40 TP-60 TP-80

T2 -54.21 -54.7 -59.86 -64.92 -

Peak localization (ppm) T3 Q2 -62.33 -90.82 -60.66 -95.03 -94.19 -72.65 -102.19 -

Q3 -99.41 -105.55 -100.74 -110.74 -101.56

Q4 -113.22 -110.16 -121.53 -

Table 2. Quantitative analysis for 29Si spectrum.

Hybrid Film TP TP-20 TP-40 TP-60 TP-80

Proportion (%) Relative proportions Dc(T) Dc(Q) %Ti (%) (%) T1 T2 T3 Q2 Q3 Q4 T1 T2 T3 Q2 Q3 Q4 3.83 8.57 10.33 16.88 40.76 19.63 16.84 37.71 45.45 21.84 52.75 25.41 76.21 75.89 22.73 14.09 11.42 10.67 42.69 11.74 9.39 38.94 31.56 29.50 66.89 18.39 14.72 63.52 61.96 36.18 24.69 18.52 0 37.65 19.14 0 57.15 42.86 0 66.30 33.70 0 47.62 58.42 43.21 11.78 12.42 6.21 41.54 21.41 6.64 38.73 40.85 20.42 59.69 30.77 9.54 60.56 62.46 30.41 0 0 0 0 100 0 0 0 0 0 100 0 0 75 0

Qi relative proportion =

Qi ∑ Qi

%Qi

TDc (%)

77.27 63.82 56.79 69.59 100

75.96 62.52 53.76 61.88 75

×100 (2)

Polymerization or condensation degree of the trifunctional %Dc(T) and tetrafunctional %Dc(Q) species was calculated using the equations (Equations 3 and 4)[22,24]: Figure 9. PEG 1500 13C NMR spectrum in liquid phase.

= % Dc (T )

T 1 + 2T 2 + 3T 3 ×100 (3) 3

= % Dc ( Q )

Q1 + 2Q 2 + 3Q3 + 4Q 4 ×100 (4) 4

The total fractions of trifunctional and tetrafunctional species were calculated as follows (Equations 5 and 6)[22,24]:

Figure 10. T and Q silicon-based species. i

350

i

= %Ti

Total T species ×100 (5) Total T species + Total Q species

%Qi =

Total Q species ×100 (6) Total T species + Total Q species Polímeros, 27(4), 346-352, 2017

Effect of concentrations of plasticizers on the sol-gel properties developed from alkoxides precursors The total degree of condensation corresponds to (Equation 7)[22,24]:

(

) (

)

1 Total % Dc = ×  % Dc (T ) × %T i + % Dc ( Q ) × %Qi  (7)  100 

The 29Si spectrum (Figure 8) shows the presence of the different silicon-based species formed in the hybrid films.

Both trifunctional (Ti) and tetrafunctional (Qi) species were found in the hybrid films, except in the one with the highest PEG concentration (TP-80). Peaks for all the silicon-based species are reported in Table 2. Peak values are similar to those found by other authors[22,23,25,26].

PEG addition diminished the formation of diverse Qi structures and favored the formation of one of them, in this case Q3 structures regarding the peak position close to -100 ppm, and the decrease or absence of other Qi peaks (Figure 10). On the other hand, PEG addition in concentration up to 60 g.L-1 increased the formation of different Ti structures. Increasing PEG concentration above that value can reduce the Ti structures amount, as the results for sample TP-80 suggest (Table 2).

The sample without PEG addition TP-0 had the lowest T1, T2 and %Ti, as well as the maximum Q3, Q4 and %Qi. It can be stated that the addition of plasticizer up to the determined concentration, in this case 40 g.L-1, enhanced the formation of a more flexible structure. In the case of absence of the plasticizer, there was a formation of a more or less condensed structure.

For all hybrid films, except for the TP-80 one, similar polymerization degrees between Ti and Qi species were observed (Table 2). However, PEG inhibited the complete reaction of all the species in the condensation reactions, since its addition enhanced the formation of T1 species[22].

Besides, the higher the amount of T1 species the higher the free volume in the polymer.

The addition of a low amount of PEG enhanced an optimal degree of polymerization for the TP-20 sample.

The formation of diverse Ti and Qi structures was promoted

by PEG addition. This allowed an optimal compromise between the compactness of Qi structures that enhances the barrier effect of the film, and the flexibility of Ti species that made the structure more resistant to plastic deformation thus avoiding cracks on the film.

Due to the PEG trend to polymerize in a semi-crystalline form, it is possible to propose that, for for TP-60 and TP-80 samples samples, PEG polymerization during the condensation reactions promoted the formation of a more fragile structure. It explains the increase of the width and the area of peaks at ppm close to 69.49 in 13C spectrum.

As well as the higher brittleness of sample TP-80.

Figure 11 shows the Fourier transform infrared (FTIR) spectra for all the hybrid films. In all the spectra, strong bands located at 1000 cm–1 and 1200 cm–1 were observed. They are attributed to Si-O-Si bonds and constitute the hybrid structure Backbone. Bands between 900 cm–1 and 960 cm–1 emerge due to the SiOCH2CH3, a product of the incomplete TEOS hydrolysis. Bands at 1728 cm–1 and 1622 cm–1 are respectively associated with C=O and C=C stretching. Peaks at 2900 cm–1 are related to symmetric and assymetric stretching of CH, Polímeros, 27(4), 346-352, 2017

Figure 11. FTIR spectra for the studied hybrid films.

CH2 and CH3 in the TMSPMA aliphatic chain. A wide band between 3200 cm–1 and 3700 cm–1 characteristic of –OH axial deformation is also present[27], which can arise from silanol groups (Si-OH) that did not condensed during the synthesis process[28]. This band increased its intensity with the augments of PEG concentration.

4. Conclusions Results showed that the increase of the T1 was observed

by RMN analysis. It promoted the formation of a compact film, enhanced by Q2 and Q3 species, with a high flexibility given by the T1 and T2 structures.

Regarding the TP-0, according to 29Si NMR spectrum, it possibly could present a good barrier effect behavior, promoted by the presence of more Qi structures. It was observed by NMR 13C spectrum and XRD diffractogram that higher PEG concentrations produced a cross-linked polymer with an important crystalline fraction.

Based on the results, it concluded that the hybrid films showed a good properties with a lower concentration of PEG. Thus, these films may be indicated for coating metal substrates aiming to improve the corrosion protection in many areas of the metals industry.

5. Acknowledgements The authors would like to thank the Brazilian government agencies CNPq and CAPES for their financial support for this research.

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Received: Oct. 12, 2016 Revised: Mar. 23, 2017 Accepted: Apr. 09, 2017

Polímeros, 27(4), 346-352, 2017

http://dx.doi.org/10.1590/0104-1428.16016

Evaluation of Out-of-Autoclave (OOA) epoxy system Fernanda Guilherme1,2, Silvana Navarro Cassu1,3, Milton Faria Diniz3, Tanila Penteado de Faria Gonzales Leal2, Natália Beck Sanches4 and Rita de Cássia Lazzarini Dutra1* Instituto Tecnológico de Aeronáutica – ITA, São José dos Campos, SP, Brazil 2 Embraer S.A., São José dos Campos, SP, Brazil 3 Divisão de Química – AQI, Instituto de Aeronáutica e Espaço – IAE, São José dos Campos, SP, Brazil 4 Centro Logístico da Aeronáutica – CELOG, Santana, SP, Brazil 1

*ritacld@ita.br

Abstract Epoxy resins (EP) usually cure in autoclave to minimize resin voids and to achieve the desired resin/fiber ratio.

Cure parameters such as temperature, vacuum and pressure levels are controlled and monitored. Aiming time and cost optimization, new out-of-autoclave (OOA) cure processes have been developed lately. This study evaluated the cure cycle and the effect of non-programmed interruptions in an OOA process. Fourier Transform Infrared spectroscopy (FT-IR) results show similarities between the resin used and diglycidyl ether of bisphenol A (DGEBA) and also that the curing system is composed of cyan and sulfur hardeners, codified in industry, as Components of #2511 Resin System.

The cure cycle and its interruptions were simulated by dynamic-mechanical analysis (DMA). The samples obtained were evaluated by FT-IR and differential scanning calorimetry (DSC), whose results show that the degree of cure varying between 0.8 to 0.85 was achieved at 120 °C. Keywords: DMA, DSC, epoxy system, FT-IR, OOA.

1. Introduction Composite materials are used in a wide range of application, such as the aerospace industry. They can be employed as rigid thermal insulation of rocket engine and in many types of aircraft parts (fairings, fuselages, leading edges)[1].

Most established curing process for high performance aerospace composites parts requires the use of autoclaves[2], which involves temperature, vacuum and pressure.

Consequently, the process present high operating costs, long cycle times and limitation of produced parts size[3], besides high investments are needed.

During the last four decades, the use of composites has undergone many changes. Several studies proposed changing the method of composite production, the most significant being the out-of-autoclave (OOA) cure[4].

Nowadays, OOA curing process is still seldom used in comparison to the conventional methods such as autoclave curing prepregs. The latest aircraft development programs, in large companies, invest millions in autoclaves and automated production machines to meet their high demands, around 80 to 120 aircraft per year. About 10% of production costs represent the curing and assembly of the manufactured composite parts. This fact is one of the motivations for the OOA process development[4].

According to Harshe[1], Thomas et al.[5] and ETH[6], prepreg (reinforcement of preimpregnated fibers with a partially cured polymer resin) of OOA is a potential alternative to materials cured in traditional autoclave.

It reduces processing costs and allows the manufacturing of larger parts, limited in conventional prepregs by the autoclave size. OOA process differs from the conventional ones because it is exclusively dependent of vacuum and

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temperature and can cure in an oven. These OOA systems are capable of achieving mechanical properties similar to some prepreg processed by autoclave. However, according to Harshe[1], further research should be developed to improve the robustness of these processes.

Epoxy resins (EP) composites are the most used in aeronautical applications due to the cost/benefit relation between price, weight, thermal and mechanical properties.

Epoxies are the resins used in OOA systems along with selected curing agents (CA)[7]. Since prepregs are supplied ready for use, their exact chemical composition is unknown and varies from manufacturer to manufacture.

There is a large amount of studies related to the characterization and identification of curing agents using Fourier Transform Infrared spectroscopy (FT-IR), by transmission mode. Some of these studies, were developed at the Brazilian Institute, Instituto de Aeronáutica e Espaço (IAE).

Sugita et al.[8] showed that carbonylated compounds are more easily identified by FT-IR due to the presence of C = O of the ester functional group, origined from the epoxide ring/anhydride curing reaction. However, the methodology used for the separation of this component is complex. A few studies reported a simpler method for CA separation: controlled pyrolysis (CONTROLPIR/FT-IR)[9,10]. These routines were able to isolate and characterize, by FT-IR, the CA anhydride type in EP and even the composite amide-amine type, which shows a higher degree of difficulty to be separated.

Romão et al.[11] demonstrated that it was possible to identify polymercaptan and amine adduct based CA, used in DGEBA, under appropriate conditions, and also that

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O O O O O O O O O O O O O O O O

Guilherme, F., Cassu, S. N., Diniz, M. F., Leal, T. P. F. G., Sanches, N. B., & Dutra, R. C. L. the developed methodology could be evaluated for the investigation of other types of CA, under specific adaptations.

Other technique coupled thermogravimetric analysis and FT-IR (TG/FT-IR) to identify polymercaptan and amine adduct based CA, as well as CA containing polyaminoamide in smaller contents.

Romão et al.[12], in another study, also developed the characterization of EP DEGBA type, cured with agents based on mercaptan and tertiary amines, by middle infrared (FT-MIR) and near infrared (FT-NIR). The spectrometric changes resulting from the curing reaction were evaluated, when the disappearance of some characteristic functional groups occurred. It was observed that the NIR region shows better the spectrometric changes of the reactions studied, allowing the detection of the CA, even in smaller proportion.

Romão et al.[13], in a supplementary study, applied the CONTROLPIR/FT-IR method to characterize DGEBA epoxy based systems cured with other agents, such as anhydride, pure or in the presence of small amounts of mercaptan and amino-phenol. The TG/FT-IR coupling was also used to identify the CA present even in a small content in the epoxy system.

Andrade et al. [14] showed that the use of CA containing mercaptan groups (CAPCURE) in DGEBA system/diethylenetriamine (DETA) reduced the cure time of an adhesive to 24 hours. FT-MIR/FT-NIR and differential scanning calorimetry analysis (DSC) results showed a complete cure. The addition of asbestos in adhesive containing CAPCURE was also evaluated. The reactions between the epoxy and amine groups, as well as those occurring between epoxy and mercaptan, were studied in the spectral regions MIR and NIR. It was observed that asbestos does not interfere in the curing reactions and the FT-NIR evidences better the spectrometric changes during reactions in comparison to the FT-MIR analysis, similar to what occurred in the study by Romão[12].

Sales et al.[15] investigated the application of luminescence spectroscopy under steady state conditions for studying F161 prepreg of fiber glass and EP. The results were compared to those in the FT-NIR analysis. Both methods indicated that the crosslinking reaction can be monitored by spectrometric analysis of the changes observed. FT-IR analysis indicated the CA used as an amine function, however it did not indicate the chemical structure of the compound. The degrees of cure evaluated by both techniques were similar.

Ferrari et al.[16] carried out the characterization of some prepregs (EP/carbon fiber (CF)) by FT-IR, TGA, DSC and dynamic-mechanical analysis (DMA). FT-IR transmission mode and photoacoustic spectroscopy (PAS) were performed.

The sample techniques used were KBr pellets, pyrolysis in Bunsen burner and controlled pyrolysis. FT-IR transmission analysis of the extract obtained by solvent treatment was able to identify cyanoguanidine (dicyandiamide) as the curing agent. FT-IR/PAS spectra evidenced the spectrometric changes resulting from the curing of the epoxy system. In addition, thermal analysis showed the curing process steps.

The application of prepreg to the manufacture of composite parts includes[16,17] cutting and laminating a predetermined number of plies to be used in a particular part. As the manufacturing process depends on a curing 354

cycle already established by the manufacturer, there is an interest of the scientific and technological community in studying which properties of the material may change once the curing process is interrupted.

In this context, the contribution of this research is the evaluation of an OOA epoxy system when compared to conventional autoclave curing system, and the influence of the curing cycle interruption at different stages of the cure.

The application of methodologies previously developed in our laboratories and the development of new methods, with the use of FT-IR latest techniques such as universal reflection attenuated (UATR) to identify OOA prepregs composition, can contribute significantly to quality control laboratories and research companies in the aeronautical industry being more competitive.

Therefore, based on the stated above and aligning the interest of companies working with prepregs, NIR and MIR techniques for the identification/characterization of the base polymer/curing agent were evaluated in this study, applied to the analysis of a epoxy system of unknown composition.

The evaluation of the materials subjected to the curing cycle interruptions at different temperatures was performed by FT-IR and DSC.

2. Materials and Methods A CF prepreg and an epoxy resin film were analyzed in this study, both of unknown composition. The analyzed prepreg has CF reinforcement, and both, prepreg and film, are made of EP of OOA curing. It was provided by Toray Company.

As prepregs are perishable materials, it was stored at -18 °C in order to slow down the chemical reactions involved in curing process and consequently keep the storage conditions during the time allowed for manufacturing parts, which is 2 years. The shelf life determined by the manufacturer is 24 months and the out time is restricted to a maximum of 25 days at room temperature[18].

2.1 Methodologies/conditions The prepreg and the EP film were previously conditioned in sealed plastic bags. For the characterization tests, they were removed from the freezer to defrost at ambient temperature for about 1 hour (23°C/65%RH). After that, the materials were cut according to required size for each technique.

2.2 FT-IR analysis They were performed on a Spectrum One PerkinElmer spectrophotometer (conditions: 7000 to 400 cm-1 for NIR analysis and 4000 to 400 cm-1 for MIR analysis, resolution 4 cm-1, 20 scans). Sample dimensions were 5 mm x 5 mm. The FT-IR spectra mode transmission were obtained directly and mode reflection by using Universal Attenuated Total Reflection (UATR) and Diffuse Reflection (DRIFT) accessories. Analyses were conducted as follows: first, non-destructive testing of prepreg was performed by MIR, UATR and DRIFT. MIR analysis was performed on the residues obtained from Sohxlet extraction. The EP film was analyzed as received by NIR. Polímeros, 27(4), 353-361, 2017

Evaluation of Out-of-Autoclave (OOA) epoxy system 2.3 Identification of CA Controlled pyrolysis was carried out at 500 °C/5 min, and then the pyrolysed film was analyzed by FT-IR transmission mode. In another step of identification, the material was treated with water in a Soxhlet apparatus, and the extract obtained was analyzed by UATR. The materials used as reference for comparison with the curing system used were curing agent and accelerator, based on cyan and amine-sulfur compounds, respectively. They also were analyzed by FT-IR mode transmission, after being prepared as KBr pellet (0.8: 400mg).

2.4 TGA analysis The materials degradation was investigated on a TGA model 4000 supplied by PerkinElmer, under the conditions: 10mg, nitrogen atmosphere (N2) (20 mL min-1), temperature range 25°C to 950°C, heating rate 10 °C/min. The purpose of this test was to determine the material degradation profile to set the optimum parameters for the pyrolysis with controlled temperature[11]..

2.5 DSC analysis The DSC analyses of the prepreg were carried out using a DSC 8000 PerkinElmer, with 10mg of samples in aluminum hermetic pans, at 10 °C/min, between -40 °C and 230 °C. The second heating was conducted under the same conditions.

Analyses were carried out with N2 atmosphere (20 mL min-1).

The final temperature used in DSC analyses was set in accordance with the initial degradation temperature determined by TGA. Samples at different temperatures (66 °C, 88 °C, 100 °C, 120 °C and 132 °C) were obtained during the cure cycle and analyzed by DSC under the conditions described above. Samples at different temperatures (66 °C, 88 °C, 100 °C, 120 °C and 132 °C) were obtained through different cure cycles as mentioned in item Curing process interruptions. DSC analyses were performed in triplicate.

The enthalpy values were ​​ corrected considering only the mass of resin, with a standard deviation of about 10%.

This correction was made by discounting the mass of the

carbon fiber, obtained from TGA analysis. The degree of cure was estimated in relation to the enthalpy obtained by the DSC curve of the uncured resin sample.

2.6 DMA analysis The cure of the prepreg was also monitored by DMA using the curing cycle recommended by the manufacturer.

The curing cycle, ranging from 25 to 132ºC is showed in Figure 1. The equipment used was a PerkinElmer DMA model 8000. Analyses were performed using dual-cantilever clamps, frequency of 1 Hz, and the dimensions of the prepreg specimens were (50 x 5 x 0.7) mm. The heating rate used between the isothermal temperatures was 1.7 ºC/min.

In the DMA test, aluminum paper was used on the prepreg tips attached to the DMA claws, avoiding the adhesion of the resin to the claw and facilitating the removal of the sample for later tests.

2.7 Curing process interruptions Simulation of curing interruptions was carried out in DMA under the same conditions described above.

The materials obtained from the interruptions at different temperatures received codes, as follows: A (25 °C to 66 °C); B (25 °C to 88 °C); C (25 °C to 100 °C); D (25 °C to 120 °C) and E (25 °C to 132 °C).

Samples were collected at the predetermined temperatures, as presented in Figure 1, and analyzed by UATR and by DSC (under the same conditions previously described, in section 2.2, to check the curing stage of each specimen).

3. Results and Discussion Initially, materials as received were characterized by surface FT-IR techniques, such as DRIFT (Figures 2A and 2B) and UATR (Figures 2C and 2D). As it is known, the depth of penetration (pathlength) of the infrared beam in the DRIFT technique is deeper than UATR[19,20], consequently DRIFT and UATR spectra can present different results. However, the sample preparation in these FT-IR reflection modes do not require complex process. The material is analyzed directly without any treatment. DRIFT spectra usually evidence

Figure 1. Representation of the prepreg curing cycle, highlighting the curing interruptions.

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Guilherme, F., Cassu, S. N., Diniz, M. F., Leal, T. P. F. G., Sanches, N. B., & Dutra, R. C. L. surface coating and inner composition showing then in some cases, large bands. Meanwhile UATR is considered a surface technique that use a zinc selenide crystal with diamond to analyze the surface material pressuring the sample against the crystal to assure that it has a good contact with the surface of sample and the incident IR beam, preventing loss of beam irradiation[21,22].

The major bands showed by both techniques, in all spectra at 3450 cm-1 (OH), 1250 and 1040 cm-1 (aromatic ether group), 917 cm-1 (main terminal epoxy band) and 826 cm-1 (para-substitution aromatic) characterized[19] the functional groups of EP. UATR spectra showed more defined resin bands, whereas DRIFT spectra distinctly showed an absorption at approximately 2160 cm-1 (Figure 2A), suggesting the detection of another component. This component might possibly be the curing agent, as the band can be assigned to cyan groups[19]. This possibility can be further investigated by other FT-IR techniques, with or without prior solvent extraction[19,23].

All data suggest so far that an effortless non-destructive analysis without samples pretreatment, by UATR and DRIFT techniques, is able to indicate the type of resin, in this case

an EP DGEBA type. It can also suggest the curing agent used as a cyan compound.

3.1 Transmission NIR/MIR analysis of the film EP To attempt a better characterization of the EP, the EP film as received was also analyzed by NIR. Figure 3 shows the NIR spectrum of the EP film. Table 1 indicates the main NIR bands and their probable assignment[12,24] which are characteristic of a DGEBA resin type.

It is known that the MIR spectra of cured epoxy resins derived from bisphenol A and epichlorohydrin are quite similar to the uncured resins spectra. The main difference is the absence of absorption at 917 cm-1, assigned to the epoxy group[25] and to the presence of new bands from the CA employed. In order to try to identify the CA, an attempt was made for controlled pyrolysis (CONTROLPIR / FT-IR), already developed by the group[9,10].

The temperature range used in the CONTROLPIR/FT-IR technique, was set in to the temperature established by TGA and corresponds to the resin degradation temperature.

Figure 2. FT-IR spectra of the analyzed materials as received: (A) OOA CF/EP prepreg (DRIFT); (B) OOA EP Film (DRIFT); (C) OOA CF/EP prepreg (UATR); (D) OOA EP Film (UATR).

Figure 3. Transmission FT-NIR partial spectrum (7600-4000 cm-1) of OOA EP film.

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Evaluation of Out-of-Autoclave (OOA) epoxy system The thermal decomposition of the film occurred in a single stage, between 290°C and 520 °C (Figure 4).

Although the TGA experiment was not able to separate the EP components, FT-IR analyses were performed on the EP film after controlled pyrolysis[9] at 500 °C, in an attempt to determine whether it would be possible to identify the CA functional groups (Figure 5C).

The bands in Figure 5C, around 3300 cm -1 and 1500 to 1600 cm-1 indicate the presence of the NH

group, and those from 1200 to 1400 cm-1, the CN bond[19].

Detection of these functional groups is in agreement with the molecular structure of cyan hardener.

However, as bands in the fingerprint region (i.e., wavenumbers under 1500 cm-1) are characteristic of specific kinds of bonds, they must be associated to another region for the unique identification of the material. In this context, the band between 2100 and 2200 cm-1 of the cyan group was not observed, suggesting that the spectrum

Table 1. Probable assignment of main NIR bands of EP film.

Wavenumber (cm-1) 6069 5984 5891 4666 and4616 4523 4360 4209 4135 4065

Probable assignment[12,24] First overtone of terminal CH2 stretching mode at 3055 cm-1 Combination band or first overtone of aromatic CH stretching at 2994 cm-1 Combination band or first overtone of aromatic CH stretching at 2966 cm-1 Combination band of aromatic C=C conjugated at 1611 cm-1 and CH stretching of epoxy ring at 3055 cm-1.

Combination band of CH2 bending or C=C conjugated at 1486 cm-1 and CH stretching of epoxy ring at 3055 cm-1.

Combination band of CH stretching of epoxy ring at 3055 cm-1 and CH3 bending at 1361 cm-1 .

Combination band of CH2 stretching at 2922 cm-1 and CH3 bending at 1361 cm-1 .

Combination band of CH2 stretching at 2922 cm-1 and aromatic ether at 1234 cm-1 Combination band of CH2 stretching at 2922 cm-1 and aromatic ether at 1148 cm-1

Figure 4. TGA and DTG curves of carbon fiber sample with epoxy resin, starting at a temperature of 30 °C to 950 °C, under nitrogen atmosphere.

Figure 5. Transmission FT-IR/MIR spectra: (A) Cyan hardener compound (KBr pellet - 0.8: 400mg), (B) extract - solid residue of EP film, obtained after treatment with water without heating (KBr pellet - 0.6: 300 mg), (C) pyrolysed EP film, obtained by controlled pyrolysis at 500 °C for 5 min.

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357

Guilherme, F., Cassu, S. N., Diniz, M. F., Leal, T. P. F. G., Sanches, N. B., & Dutra, R. C. L. (Figure 5C) shows only one pyrolysis product that contain amine groups. Results, however, can not indicate undoubtedly that it is a known cyan hardener, as Figure 5A, for example.

Indication of cyan groups in the prepreg and resin let to the assessment of various types of solvents, as previously performed by Ferrari et al.[16], for solvent treatment, which is done to isolate and study the different components.

The best result for the separation and identification of the CA was observed by using water as solvent. This is accordance with the literature[23], that indicates that cyan hardener compound, containing NH2, NH and CN groups, is partially soluble in water.

The comparison between Figures 5A and 5B indicates that the CA used is based on a cyan compound[26]. Therefore, as the spectrum of the controlled pyrolysis (Figure 5C) failed to present any cyan functional groups, it can be concluded that

the water extraction was more effective for the identification of cyan hardener, probably due to higher concentration of this compound in extract than in pyrolysis product of prepreg, which contains traces of compound.

3.2 FT-IR analysis of the prepreg after solvent treatment Literature[27] states that the cure of epoxy systems with cyan hardener compound often requires the presence of a additional compound in the formulation, such as an accelerator. Its function is to work as the system activator.

The prepreg composition is not completely known and it is reasonable that another CA or additives were added to the formulation along with the cyan hardener compound.

For this reason, the prepreg was treated with pyridine[23], which is a solvent currently used for extracting sulfur curing accelerator in prepregs at the Brazilian Institute, IAE. Figure 6 shows the FT-IR spectra obtained after extraction with pyridine and previously treated with acetone. Absorptions at 3459 and 3367 cm-1 can be attributed[19] to NH groups, those at 1285, 1144, 1105 and 600 cm-1 suggest the presence of sulfone groups, and the one at 3060 cm-1 is likely due to the CH aromatic group. The doublet at 1593-1502 cm-1 can be assigned to the C-C aromatic group, while the band at 830 cm-1 assigned to the para-substituted CH aromatic.

Figure 6 shows also the reference spectrum for an amine-sulfur curing accelerator[26]. The similarities observed indicate this compound is in the curing system. Hence, it can be concluded that the epoxy resin is of the DGEBA type containing cyan and sulfur compounds, codified in industry, as hardener components of #2511 system.

3.3 Evaluation of resin curing behavior Figure 6. Transmission FT-MIR spectra of: (A) Extract- residue obtained after treatment with acetone and pyridine (KBr pellet 0.8: 400 mg); (B) amine-sulfur curing accelerator (cast) (KBr pellet- 0.8: 400 mg KBr).

The resin curing cycle as recommended by the manufacturer was simulated by DMA (Figure 7). Initially, as temperature increases, the storage modulus (E’) profile shows a decrease. This happens prior to the curing reaction,

Figure 7. Initial characterization of Prepreg DMA curves: storage modulus curve (__), loss modulus (---) and tan δ (-.-.-) of CF/ EP obtained in heating rate 1.7 °C/min in temperature. Representation of the prepreg curing cycle, highlighting curing interruptions in their respective temperature ranges: A (25 °C to 66 °C); B (25 °C to 88 °C); C (25 °C to 100 °C); D (25 °C to 120 °C) and E (25 °C to 132 °C).

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Evaluation of Out-of-Autoclave (OOA) epoxy system due to the reduction of resin viscosity as it is heating up.

This behavior contributes to a better spread of the resin into the prepreg fibers.

After 150 min, it was noted that the storage modulus increased gradually. It indicates that the resin is gelling, what decreases the polymeric matrix flowability[18]. As time goes by, E’ increases continuously until it reaches a stable value around 190 min what indicates that the cure reactions are completed at 132°C. The curing behavior was also monitored by DSC, at constant heating rate of 10 °C/min. The DSC curve (Figure 8) of the CF/EP prepreg reach its maximum exothermic peak at 158 °C. The enthalpy regarding the curing reaction of the epoxy resin is 164 J/g. In Figure 8, the partial Tg of the system can be observed close to 0 °C of the uncured prepreg sample and this Tg disappears with the thermal cycles shown in Figure 9.

Figure 8. Prepreg DSC curves of CF/EP at 10 °C/min in N2 atmosphere.

Aeronautical parts manufactured from OOA systems can have its curing cycles unintentionally interrupted for several reasons. Interruptions of curing cycles at different temperatures can be simulated in DMA and consequences in their chemical structure can be evaluated by FT-IR and DSC. DMA analysis initially made with prepreg in constant heating rate was used as a reference for choosing the temperatures in which to interrupt cycles, in accordance to the variation of storage modulus (as shown in Figure 7).

The temperature chosen for the first curing interruption was 66 °C (sample A), because the material at this point had not undergone a significant change in its composition, only variation of resin viscosity without changes in storage modulus (E’). An interruption at this temperature probably does not affect the material curing process. DSC (Figure 9) and FT-IR (Figure 10) analysis on sample A shows that there was practically no variation from the raw material results.

Sample B was obtained after the second interruption of the curing cycle at 88 °C. FT-IR and DSC results of sample B also showed a similar profile to the prepreg prior to curing process. This behavior was expected due to the storage modulus curve, which only presents an increase after 140 minutes after starting the analysis.

Storage modulus of samples C and D, from interruptions at 100 °C (140 min) and 120 °C (155 min), respectively, showed an increase in value. Sample C showed a slight reduction in the residual cure in comparison with samples A and B, showing the degree of cure of 0.3. Sample D results indicate that the curing process is almost complete. It can be evid.enced by the observation of the DSC curves, as samples D and E reached a degree of cure of 0.8 and 0.85, respectively. FT-IR analysis of samples D and E (Figure 10 - D and E) showed that the band assigned to the epoxy group (around 910 cm-1) is not observed at 120 °C. These findings are in accordance to the decrease of curing reaction accompanied by DSC. This curing temperature is coincident with the one cited by Ferrari et al.[16] for a prepreg used in the aeronautical industry.

In temperatures higher than 120°C, bands of CO/CN in the region between 1000-1200 cm-1, and of aromatic substitution around 800 cm-1, suffer modifications in position, shape and intensity, suggesting changes in chemical structures[19].

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Figure 9. DSC curves of CF prepreg with EP obtained after isothermal treatment (curing interruptions). Delta H was corrected in relation to the mass of resin.

Figure 10. UATR spectra of the OOA epoxy system: all samples will start at 25 oC: (A) until 66 °C; (B) 88 °C; (C) until 100 °C; (D) until 120 °C; (E) until 132 oC, (both at a heating rate of 1.7 °C/min).

4. Conclusion FT-IR analysis of OOA epoxy system in the form of prepreg and film showed that the resin has a structure similar to the one found in DGEBA-based polymers and that the curing system used is based on cyan hardener and 359

Guilherme, F., Cassu, S. N., Diniz, M. F., Leal, T. P. F. G., Sanches, N. B., & Dutra, R. C. L. amine-sulfur compounds, codified by industry as hardener components of #2511 system.

The DSC evaluation of interruptions in the curing cycle at different temperatures, showed that interruptions in temperatures up to 100 °C should not interfere with the continuity of the cure process, since it was observed a low degree of cure (0.3 at 100 °C). The FT-IR spectra of these samples detected the same characteristic bands as the FT-IR spectrum of the resin prior to the cure process, evidencing minor or no changes whatsoever in the chemical composition.

On the other hand, DSC results of samples collected from interruptions at temperatures higher than 120 °C showed that the resin was almost completely cured, with a degree of cure between 0.8 and 0.85. Furthermore, the FT-IR spectra of these samples did not evidenced the band related to the epoxy functional group, validating the DSC findings.

Based on the results found, we consider that the methodology developed in our laboratories, by using of these spectroscopic techniques, FT-MIR (transmission and reflection modes, including UATR) and FT-NIR (transmission mode), associated to thermal analyses such as TGA, DSC and DMA, can be an important contribution for studying of epoxy systems composition and evaluation of OOA curing behavior.

5. Acknowledgements This paper was supported in part by PVNS (National Senior Visiting Professor Program) and PROAP (Postgraduate Support Program) from CAPES and ADC/DCTA (Classical sports association of civil and military servants of the Aerospace Technical Center).

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17. Pardini, L. C., & Peres, R. J. C. (1996). Tecnologia de fabricação de pré-impregnados para compósitos estruturais utilizados na indústria aeronáutica. Polímeros: Ciência e Tecnologia, 6(2), 32-42. Retrieved in 2015, June 25, from http://www. revistapolimeros.org.br/PDF/v6n2/v6n2a02.pdf 18. Hexcel. (2017). Prepreg technology. Stamford: Hexcel. Retrieved in 2017, February 22, from https://www.ethz.ch/content/ dam/ethz/special-interest/mavt/design-materials-fabrication/ composite-materials-dam/Education/Manufacturing_of_ Polymer_Composites/FS2017/Prepreg_Technology.pdf 19. Smith, A. L. (1979). Applied infrared spectroscopy: fundamentals, techniques, and analytical problem-solving. New York: John Wiley & Sons.

20. Graf, R. T., Koenig, J. L., & Ishida, H. (1987). Fourier transform infrared characterization of polymers. New York: Plenum Press.

21. Arrizabalaga, I., Gomez, L. O., Aramendia, J., Arana, G., & Madariaga, J. M. (2014). Determination of the pigments present in a wallpaper of the middle nineteenth century: The combination of mid-diffuse reflectance and far infrared Polímeros, 27(4), 353-361, 2017

Evaluation of Out-of-Autoclave (OOA) epoxy system spectroscopies. Spectrochimica Acta Part A: Molecular and Biomolecular Spectroscopy, 124, 308-314. http://dx.doi. org/10.1016/j.saa.2014.01.017.

22. Noureddine, A., Hequet, E., Turner, C., & Sari-Sarraf, H. (2005). FTIR analysis of crosslinked cotton fabric using a ZnSeâ&#x20AC;&#x201C;universal attenuated total reflectance. Journal of Applied Polymer Science, 96(2), 392-399. http://dx.doi.org/10.1002/ app.21449.

23. Hawley, G. G. (1981). The condensed chemical dictionary (10th ed.). New York: Van Nostrand Reinhold Company.

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25. Urbanski, J., Czerwinski, W., Janicka, K., Majewska, F., & Zowall, H. (1977). Uncured epoxy resin: handbook of analysis of sinthetic polymers and plastics. New York: John Wiley & Sons.

26. Hummel, D. O., & School, F. (1981). Spectra and methods of identification: atlas of polymer and plastics analysis (Vol. 3).

Deerfield Beach: Dr. Hans F. Ebel.

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New York: McGraw-Hill.

Received: Nov. 22, 2016 Revised: Mar. 07, 2017 Accepted: Apr. 22, 2017

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40.00

35.00

APC vs. GPC com padrão de poliestireno Mp = 510 30.00

µRIU

25.00 40.00 20.00 35.00 15.00 30.00

µRIU

10.00 25.00 5.00 20.00 0.00 15.00 1.60

1.80

2.00

2.20

2.40

2.60

10.00

APC 2.80

3.00

3.20

3.40

3.60

3.80

4.00

3.00

3.20

3.40

3.60

3.80

4.00

Minutes

5.00

0.00 1.60

1.80

2.00

2.20

2.40

2.60

2.80

Minutes 22.0

gPC

20.0 18.0 16.0

µRIU

14.0 22.0 12.0 20.0 10.0 18.0 8.0 16.0

µRIU

6.0 14.0 4.0 12.0 2.0 10.0 0.0 8.0 4.00

4.20

4.40

4.60

4.80

5.00

5.20

5.40

5.60

5.80

6.00

6.20

6.40

6.0

Minutes

4.0 2.0 0.0

O CAMINHO PARA A INOVAÇÃO. 4.00

4.20

4.40

4.60

4.80

5.00

5.20

5.40

5.60

5.80

6.00

6.20

6.40

Minutes

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