Polímeros: Ciência e Tecnologia (Polimeros), vol.25, n.3, 2015

Polímeros

VOLUME XXV - N° 3 - MAIO/JUN - 2015

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Editorial

http://dx.doi.org/10.1590/0104-1428.2503

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Elisabete Frollini

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Agradecemos a todos (FAPESP, CNPq, empresas), e contamos com a continuidade deste inestimável apoio, o qual é imprescindível para a manutenção das mencionadas conquistas, assim como para o avanço na direção de um periódico cada vez mais em consonância com outros consolidados internacionalmente.

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As organizações sem fins lucrativos têm desempenhado um destacado papel na divulgação de resultados científicos, como é o caso da Associação Brasileira de Polímeros, que publica este periódico. A manutenção da periodicidade, atualmente bimestral, a introdução de modificações recentes, como a publicação em XML (eXtensible Markup Language), a introdução de um Template para elaboração dos manuscritos, a adoção do inglês como linguagem em todo o sistema ligado a esta revista, dentre outros fatos, contou com o importante apoio da Fundação de Amparo a Pesquisa do Estado de São Paulo (FAPESP), do Conselho Nacional de Desenvolvimento Científico e Tecnológico (CNPq), e das empresas que apóiam Polímeros.

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Neste contexto, tem crescido progressivamente o número de artigos redigidos nesta língua que são publicados em Polímeros: Ciência e Tecnologia. No presente fascículo, cerca de 80% dos artigos foram redigidos em inglês, sendo nosso objetivo que em futuro próximo esta seja a língua oficial deste periódico. Solicitamos especial atenção dos leitores para este fato, e sugerimos fortemente que compartilhem conosco deste objetivo, no momento em que decidirem submeter um manuscrito para Polímeros.

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Comunicação é uma questão central da atualidade, e a informação é fundamental na área científica. Os cientistas desejam comunicar os resultados das respectivas pesquisas para a comunidade global, o que só é possível se uma linguagem internacionalmente conhecida for usada, ou seja, se os textos científicos forem redigidos em inglês.

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Polímeros, 25(3), 2015

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P o l í m e r o s - N º 3 - V o l u m e X X V - M a i o / J u n - 2 0 1 5 - ISS N 0 1 0 4 - 1 4 2 8

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I n d e x a d a : “ 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 ” — “ I S I W e b o f K n o w l e d g e , W e b o f S ci e n c e ”

and

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Polímeros P r e s i d en t e

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Conselho Editorial

Comitê Editorial

Marco-Aurelio De Paoli (UNICAMP/IQ)

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Membros

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Conselho Editorial

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Adhemar C. Ruvolo Filho (UFSCar/DQ) Ailton S. Gomes (UFRJ/IMA) 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) Elias Hage Jr. (UFSCar/DEMa) Elisabete Frollini (USP/IQSC) Eloisa B. Mano (UFRJ/IMA) Glaura Goulart Silva (UFMG/DQ) 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) Osvaldo N. Oliveira Jr. (USP/IFSC) Raquel S. Mauler (UFRGS/IQ) Regina Célia R. Nunes (UFRJ/IMA) Rodrigo Lambert Oréfice (UFMG/DEMET) Sebastião V. Canevarolo Jr. (UFSCar/DEMa) Silvio Manrich (UFSCar/DEMa)

Elisabete Frollini – Editora

Membros

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Comitê Editorial

Adhemar C. Ruvolo Filho Bluma G. Soares César Liberato Petzhold Glaura Goulart Silva José António C. Gomes Covas José Carlos C. S. Pinto Regina Célia R. Nunes Sebastião V. Canevarolo Jr.

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Produção

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Assessoria Editorial

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www.editoracubo.com.br

Tiragem 1.500 exemplares

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“Polímeros” é uma publicação da Associação Brasileira de Polímeros Rua São Paulo, nº 994 CP 490 - 13560-340 - São Carlos, SP, Brasil Fone/Fax: (16) 3374-3949 e-mails: abpol@abpol.org.br / revista@abpol.org.br http://www.abpol.org.br

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Data de publicação: Junho de 2015 Apoio:

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Polímeros / Associação Brasileira de Polímeros. vol. 1, nº 1 (1991) -.- São Carlos: ABPol, 1991Versão eletrônica disponível no site: www.scielo.br

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Bimestral v. 25, nº 3 (maio/jun. 2015) ISSN 0104-1428

Site da Revista “Polímeros”: www.revistapolimeros.org.br

1. Polímeros. l. Associação Brasileira de Polímeros.

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Polímeros, 25(3), 2015

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Polímeros Seção Editorial

S e ç ã o T é cn i c a

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Informes & Notícias ............................................................................................................................................................................E4 Calendário de Eventos ........................................................................................................................................................................E5 Associados...........................................................................................................................................................................................E6

Development of dual-sensitive smart polymers by grafting chitosan with poly (N-isopropylacrylamide): an overview Nívia do Nascimento Marques, Ana Maria da Silva Maia and Rosangela de Carvalho Balaban.................................................................. 237

TG/FT-IR characterization of additives typically employed in EPDM formulations Blends of ground tire rubber devulcanized by microwaves/HDPE - Part A: influence of devulcanization process

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Natália Beck Sanches, Silvana Navarro Cassu and Rita de Cássia Lazzarini Dutra...................................................................................... 247 Fabiula Danielli Bastos de Sousa, Júlia Rocha Gouveia, Pedro Mario Franco de Camargo Filho, Suel Eric Vidotti, Carlos Henrique Scuracchio, Leice Gonçalves Amurin and Ticiane Sanches Valera..................................................................................... 256

Immobilization of myoglobin in sodium alginate composite membranes Katia Cecília de Souza Figueiredo, Wilbert van de Ven, Matthias Wessling, Tito Lívio Moitinho Alves and Cristiano Piacsek Borges........ 265 Tayfun Uygunoglu, Witold Brostow and Ibrahim Gunes.................................................................................................................................. 271

Analysis of equations of state for polymers

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Wear and friction of composites of an epoxy with boron containing wastes

Erlí José Padilha Júnior, Rafael de Pelegrini Soares and Nilo Sérgio Medeiros Cardozo.............................................................................. 277

Acacia bark residues as filler in polypropylene composites

Melting and crystallization of poly(3-hydroxybutyrate): effect of heating/cooling rates on phase transformation Renate Maria Ramos Wellen, Marcelo Silveira Rabello, Inaldo Cesar Araujo Júnior, Guilhermino José Macedo Fechine and Eduardo Luis Canedo............................................................................................................................................................................... 296

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Ticiane Taflick, Élida Gonçalves Maich, Laís Dias Ferreira, Clara Isméria Damiani Bica, Silvia Rosane Santos Rodrigues and Sônia Marlí Bohrz Nachtigall................................................................................................................................................................... 289

Development of paints with infrared radiation reflective properties Eliane Coser, Vicente Froes Moritz, Arno Krenzinger and Carlos Arthur Ferreira........................................................................................ 305

Bioactivity, biocompatibility and antimicrobial properties of a chitosan-mineral composite for periodontal tissue regeneration Desenvolvimento da metodologia para síntese do poli(ácido lático-co-ácido glicólico) para utilização na produção de fontes radioativas

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Andrew Paul Hurt, Arun Kumar Kotha, Vivek Trivedi and Nichola Jayne Coleman....................................................................................... 311 Fernando dos Santos Peleias Junior, Carlos Alberto Zeituni, Maria Elisa Chuery Martins Rostelato, Guilhermino José Macêdo Fechine, Carla Daruich de Souza, Fábio Rodrigues de Mattos, Eduardo Santana de Moura, João Augusto Moura, Marcos Antônio Gimenes Benega, Anselmo Feher, Osvaldo Luiz da Costa e Bruna Teiga Rodrigues.................................................................................................................. 317

Laminados biodegradáveis de blendas de amido de mandioca e poli(vinil álcool): efeito da formulação sobre a cor e opacidade

Capa: Foto - Tg values of the devulcanized rubber and rubber phases of the blends as determined by DMA. G’ and η* versus frequency of the blends. The curves were separated for better visualization and analysis of results. SEM micrographs of the blends: (a) 80GTR0/20HDPE, (b) 80GTR0+ad/20HDPE.

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Elaboração artística Editora Cubo.

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Juliano Zanela, Mônica Oliveira Reis, Adriana Passos Dias, Suzana Mali, Maria Victória Eiras Grossmann e Fabio Yamashita............... 326

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I N F O R M E S

BIO-BASED PRODUCTS: EU-funded “Placard project” reports good test results for PVC plasticiser based on cashew nut shells The latest tests conducted within the framework of the “Placard project” (www.placard-ecoinnovation. eu), which is developing a new bio-based plasticiser for soft PVC, have shown promising results, the project’s partners say. The EU-funded initiative is using cardanol – a yellow oil obtained by the vacuum distillation of cashew nut shells – as a base for the plasticiser, and pilot scale production of the material is being carried out at the premises of project member Serichim (Torviscosa, Udine / Italy; www.serichim. it), a contract research organisation.

The material has been tested at the labs of fellow project partner Università del Salento, which compared its performance to that of established phthalate and non-phthalate plasticisers, of both high and low molecular weight. The tests showed that Placard’s material has a better plasticising efficiency than DEHP and DOTP and also showed better processability than “selected commercial plasticisers”. Energy costs were also reduced, project members said, adding that all other tested parameters were comparable to those of other selected plasticisers and “showed good stability of properties over time.”

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Other results of Placard, including the ability to facilitate recycling, will be further investigated.

The remaining two members of the 32-month project, which was set up last year, are polymer conversion specialist Kommi and European Plastics Converters.

Source: Plasteurope.com

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Braskem anuncia a construção de dois novos laboratórios para análises químicas e testes de polímeros A Braskem, maior petroquímica das Américas e líder mundial na produção de biopolímeros, anuncia a construção de dois novos laboratórios para análises químicas e testes de polímeros. O investimento de R$ 1,5 milhão tem como objetivo de aperfeiçoar o apoio no desenvolvimento de produtos e os serviços oferecidos a clientes da petroquímica. A previsão é que o os espaços, localizados no Centro de Tecnologia e Inovação, em Triunfo (RS), sejam inaugurados em outubro deste ano.

Um dos laboratórios terá como foco análises de raios-X, para avaliar conteúdos residuais e propriedades de cristalinidade dos polímeros.

A geração de conhecimento trazida por essas técnicas permitirá a criação de novos produtos com propriedades diferenciadas, bem como apoiar o desenvolvimento de catalisadores, reduzindo a dependência tecnológica da Braskem.

Já o outro espaço será dedicado à análise de barreiras à passagem de gases em películas de polímeros. A partir de um processo de simulação identifica-se a eficácia do produto para impedir a entrada de oxigênio ou outros gases em aplicações de embalagens alimentícias, farmacêuticas, bombonas agroquímicas e tanques de combustível. Dessa forma, é possível ampliar a segurança, a durabilidade e o E4

atendimento a requisitos normativos das aplicações.

A reforma implica em maior capacidade de execução dos testes já realizados pela petroquímica.

“As novidades demonstram o nosso empenho em oferecer todo o suporte necessário aos nossos clientes, ao disponibilizar estrutura para testes, medições e análises. Isso dinamiza os processos para possíveis ajustes e garante melhor qualidade do produto que chega ao mercado. Por fim, o investimento reforça o nosso compromisso com a inovação e desenvolvimento da cadeia do plástico”, afirma Nércio Hexsel, coordenador do CTI. Source: Braskem Estradas de plástico apontam futuro viável da mobilidade A tecnologia pode ser aplicada em estradas ou em área urbana e apresenta muitas vantagens sobre as formas atuais de pavimentação, como a sua estrutura tubular para a instalação de tubos e cabos, menores frequência e custo de manutenção, maior vida útil, aproveita plástico reciclado em sua composição.

A cidade portuária de Roterdã, na Holanda, pode ser a primeira do planeta a construir ruas e estradas com plástico reciclado. A empresa VolkerWessels, maior construtora para obras de infraestrutura dos Países Baixos, apresentou a tecnologia na semana passada. Trata-se de uma alternativa mais “verde” de pavimentação, considerando que o asfalto é responsável pela emissão de 1,6 milhões de toneladas de dióxido de carbono anualmente, ou 2% de todas as emissões do transporte rodoviário.

Além do imenso ganho ambiental, as vias de plástico podem ser implantadas em semanas, apenas, reduzindo drasticamente o período de obras para a construção ou reforma das rodovias, avenidas e ruas. No entanto, os benefícios ainda vão além, ao considerar que a tecnologia de plástico oferece uma vida útil até três vezes superior em relação às estradas tradicionais, segundo a empresa.

Dessa forma, o índice de manutenção e os gastos com esse serviço também despencam. As estradas de plástico serão mais leves e terão o seu interior oco, permitindo a instalação de tubulações e cabeamento de serviços públicos sob sua superfície.

A tecnologia parece tão interessante, que a pergunta mais pertinente é: quando o produto será lançado comercialmente? A empresa VolkerWessels diz que o projeto ainda está em fase conceitual, porém, pretende entregar a primeira estrada totalmente de plástico reciclado em três anos, na cidade holandesa de Roterdã, que já manifestou interesse. “O plastic road oferece várias vantagens em relação a estruturas rodoviárias em uso atualmente, tanto na construção, quanto na manutenção de estradas”, afirmou Rolf Mars, diretor de Infraestrutura da VolkerWessels.

“O potencial do conceito é enorme. Neste momento estamos à procura de parceiros para desenvolver o nosso piloto. Além de fabricante da indústria de plásticos, estamos pensando em indústria de reciclagem, universidades e outras instituições de pesquisa. Estamos muito otimistas sobre o desenvolvimento em torno do Plastic Road“, disse Jaap Peters, engenheiro da prefeitura de Roterdã.

Source: Transpo Online Polímeros, 25(3), 2015

October 2015 5th International Conference on Biodegradable and Biobased Polymers (BIOPOL-2015) Date: 6–9 October 2015 Local: Donostia - San Sebastian Website: http://www.biopol-conf.org/ European Conference on Metal Organic Frameworks and Porous Polymers 2015 Date: 12–14 October 2015 Local: Potsdam - Germany Website: http://events.dechema.de/en/euromof2015.html Painel Nordeste - Tecnologias em Composites, Poliuretano e Plásticos de Engenharia Date: 20 October 2015 Local: Recife - PE Website: http://www.tecnologiademateriais.com.br/paineis.html 11th International Conference on Advanced Polymers via Macromolecular Engineering (APME-2015) Date: 18-22 October 2015 Local: Yokohama - Japan Website: http://www.apme2015.jp/ 13º CBPol - Congresso Brasileiro de Polímeros Date: 18–22 October 2015 Local: Natal - RN Website: http://www.cbpol.com.br/ Polyolefin Additives 2015 Date: 20–22 October 2015 Local: Cologne - Germany Website: http://www.amiplastics.com/events/event?Code=C671

November 2015 Polymer Foam 2015 Date: 2-4 November 2015 Local: Cologne - Germany Website: http://www.amiplastics.com/events/event?Code=C661 Painel Espumas Flexíveis Date: 4 November 2015 Local: São Paulo - SP Website: http://www.tecnologiademateriais.com.br/paineis.html Europack - Euromanut 2015 Date: 17-19 November 2015 Local: Lyon - France Website: http://www.europack-euromanut-cfia.com/en

December 2015 Fire Resistance in Plastics 2015 Date: 14-16 December 2015 Local: Cologne - Germany Website: http://www.amiplastics.com/events/event?Code=C673 Pacifichem 2015 Date: 8-10 December 2015 Local: Hawaii - USA Website: http://www.pacifichem.org/

January 2016 Plastics in Automotive Date: 14 January 2016 Local: Detroit - USA Website: www.plasticsnews.com/2015auto

Thermoplastic Concentrates 2016 Date: 26-28 January 2016 Local: Florida - USA Website: http://www.amiplastics.com/events/event?Code=C697

February 2016 Polymers in Photovoltaics 2016 Date: 2-3 February 2016 Local: Düsseldorf - Germany Website: http://www.amiplastics.com/events/event?Code=C703 6th Annual Next Generation Bio-Based & Sustainable Chemicals Date: 3-5 February 2016 Local: New Orleans - USA Website: http://www.infocastinc.com/events/biobased-chemicals Plastec West 2016 Date: 9-11 February 2016 Local: California - USA Website: http://plastecwest.plasticstoday.com/ SPE International Polyolefins Conference Date: 22-25 February 2016 Local: Texas - USA Website: https://www.spe-stx.org/conference.php

March 2016 Sustainable Plastics 2016 Date: 1-2 March 2016 Local: Cologne - Germany Website: http://www.amiplastics.com/events/event?Code=C706 KOPLAS 2015 Date: 10-14 March 2016 Local: Goyang - Korea Website: http://www.koplas.com/ Plastimagen 2016 Date: 8-11March 2016 Local: Ciudad de México - México Website: http://www.plastimagen.com.mx/en

April PlastShow 2016 Date: 12-15 April 2016 Local: São Paulo - SP Website: http://www.arandanet.com.br/eventos2016/plastshow

May 2016 International Workshop on Polymer Reaction Engineering Date: 17–20 May 2016 Local: Hamburg - Germany Website: http://events.dechema.de/events/en/pre2016.html 26th Annual Conference on Recent Advances in Flame Retardancy of Polymeric Materials Date: 17–20 May 2016 Local: Connecticut - USA Website: www.bccresearch.com/conference/flame

July 2016 80th Prague Meeting on Macromolecules - Self-Organizaion in the World of Polymers Date: 10–14 July 2016 Local: Prague - Czech Republic Website: http://www.imc.cas.cz/sympo/80pmm/

Polímeros, 25(3), 2015 E5

Associados da ABPol Patrocinadores

Instituições 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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Polímeros, 25(3), 2015

Associados da ABPol As nossas boas vindas...

Coletivos 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.

Fastplas Automotive Ltda.

Formax Quimiplan Componentes para Calçados Ltda.

Fundação CPqD - Centro de Pesquisa e Desenvolvimento em Telecomunicações Imp. e Export. de Medidores Polimate Ltda.

Innova S/A Instituto de Aeronáutica e Espaço/AQI Jaguar Ind. e Com. de Plásticos Ltda Johnson & Johnson do Brasil Ind. Com. Prod. para Saúde 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

Polímeros, 25(3), 2015 E7

Extrusora Dupla Rosca - AX 16 DR

AX 16 Granulação

Mini Injetora - AX 16 III

Multifilamentos - AX 16 MF

AX 16 Filme Tubular - Balão

AX 16 Laminadora

AX 16 Filmes Planos (Chill Roll)

R. 23 de julho, 165 - Jd. Canhema - Diadema - SP - CEP: 09941-610 axplasticos@axplasticos.com.br - www.axplasticos.com.br

fone: 55 11 4072-1161

http://dx.doi.org/10.1590/0104-1428.1744

Development of dual-sensitive smart polymers by grafting chitosan with poly (N-isopropylacrylamide): an overview Nívia do Nascimento Marques1, Ana Maria da Silva Maia1 and Rosangela de Carvalho Balaban1 1 Laboratório de Pesquisa em Petróleo, Instituto de Química, Universidade Federal do Rio Grande do Norte – UFRN, Natal, RN, Brasil

*anamaia@mail.uft.edu.br

Abstract A great deal of research on polymers over the past two decades has been focused on the development of stimuli-responsive polymers to obtain materials able to respond to specific surroundings. In this paper, an overview is presented of the concepts, behavior and applicability of these “smart polymers”. Polymers that are temperature- or pH-sensitive are discussed in detail, including the response mechanisms and types of macromolecules, because they are easy to handle and have a wide range of applications. Finally, the combination of pH and temperature responsive properties by means of graft copolymerization of chitosan with poly (N-isopropylacrylamide) (PNIPAM) was chosen to represent some synthetic routes and properties of dual-sensitive polymeric systems developed currently.

Keywords: smart polymer, thermosensitive, pH-responsive, N-isopropylacrylamide, chitosan.

1. Introduction Recently, scientists all over the world have been attempting to synthesize polymers capable of mimic the stimuli-responsive property present in common biopolymers of living organisms, in order to reach scientific and industrial applications. Frequently named as smart polymers, those materials are able to undergo fast, abrupt and reversibly alteration in their structure/properties as a response to small changes in the environment[1-4]. This behavior can be employed in a wide range of applications, such as controlled drug delivery[5-8], chromatographic separation[9,10], water remediation[11-13], enhanced oil recovery[14,15], catalysis[16,17], sensors[18,19] and tissue engineering[20,21].

The stimuli can be classified as physical, such as temperature[22-24], electric/magnetic fields[25,26] and light[27]; chemical, such as pH[28,29] and ionic strength[30]; and biochemical, such as enzymes[31] and antigens[32]. These unique macromolecules are also known as intelligent polymers[33], responsive polymers[34], sensitive polymers[35], stimuli-sensitive polymers[36], stimuli-responsive polymers[37], environmentally sensitive polymers[38] and environmentally responsive polymers[39,40].

These materials have been designed in many forms, depending on the desired application. They can be polymers chains dissolved in solutions, chemically crosslinked hydrogels, physical gels, micelles and even chains immobilized or grafted onto solid surfaces[2,4,10]. Furthermore, regardless the physical form, the smart polymers can be conjugated with biomolecules or synthetic substances, and the activity of the conjugates is going to depend on the polymer-conjugate interactions and the response of the polymer to the stimulus applied[2,4].

In order to improve and increase their applicability, several smart polymers have been developed to combine two or more mechanisms of responsiveness in only one

Polímeros, 25(3), 237-246, 2015

material[41,42]. A quite common approach involves preparing polymers sensible to both temperature and pH stimuli[43-55], since they are typical variables parameters in biological and chemical systems, and they can also be easily controlled in vitro and in vivo conditions[56-59].

The aim of this paper is to present the methods of obtainment and properties of smart polymers in a compact form. It is not intended to be a complete review and therefore the selection of cited literature is of some extent personal.

Our presentation is focused mainly on copolymers of chitosan with poly (N-isopropylacrylamide) (PNIPAM), due to the large number of studies found in the literature on these copolymers, but many properties are universal and can also be applied to other pH and temperature sensitive polymers.

2. Thermosensitive Polymers Temperature is one of the most interesting properties investigated in responsive polymer systems[2,24,59-65].

The thermosensitive polymers are well-known by possessing a large alteration in their structure as a response to slight changes in temperature. When a polymer is dissolved in an appropriate solvent, it may become insoluble upon increase or decrease in temperature and, thus, precipitate from the solution[66]. Solubility of most polymers increases with increasing temperature[59,63]. When phase separation happens with decreasing temperature, polymer system presents an upper critical solution temperature (UCST)[67]. However, some polymers exhibit a peculiar behavior, in which a phase separation occurs with rising temperature. The temperature in which this occurs is called lower critical solution temperature (LCST)[4,10,59]. The LCST behavior in water has attracted much attention due to the great applicability[66], such as in wastewater treatment[68], chromatographic separation[69], enzyme immobilization[70] and tissue engineering[71].

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Marques, N. N., Maia, A. M. S., & Balaban, R. C. From a thermodynamic point of view, the solubilization at low temperatures occurs due to the polymer-solvent hydrogen bonding that leads to a negative enthalpy of mixing. On the other hand, even with a moderate gain in compositional entropy as a consequence of the mixing process, the entropy of organization required to achieve this polymer-solvent hydrogen bonding is unfavorable (negative entropy). Thus, the free energy of dissolution, ΔG, given by ΔH - TΔS, can change from negative (solution) to positive (phase transition) as the temperature is increased[59,66,72].

The N-alkyl-substituted polyacrylamides are an especially relevant family of thermosensitive polymers, as seen by a huge numbers of publications. Interestingly, the LCST in this group of polymers varies according to the type of N-substituted groups, ranging from materials insoluble at room temperature to other with high LCST[65]. They include poly (N-isopropylacrylamide)[2,56,62,65,73-75], poly (N,N‑diethylacrylamide) [2,56,59,62] , poly (2-carboxyisopropylacrylamide), poly (N-(L)‑(1‑hydroxymethyl) propylmethacrylamide, poly (N-acryloyl-N’-propilpiperazine) [2] , poly (N-ethylacrylamide), poly (N-methyl-N-ethylacrylamide), poly (N-n‑propylacrylamide)[56,65], poly (N-ethylmethacrylamide), poly (N-methyl‑N‑isopropylacrylamide), poly (N-isopropylmethacrylamide), poly (N-npropylmethacrylamide), poly (N-methyl-Nisopropylacrylamide), poly (N-cyclopropylacrylamide), poly (N-cyclopropymethacrylamide)[56], poly (N,N-bis(2methoxyethyl) acrylamide), poly (N-(3‑methoxypropyl) acrylamide), poly (ethoxypropylacrylamide)[65] and poly (aminomethoxypropylacrylamide)[76].

Thermosensitive polymers that do not contain acrylamide‑based repeat units have also attracted significant attention, such as poly (N-vinylcaprolactam)[59,62,65], poly (2-ethyl-2-oxazoline)[65], poly (vinyl methyl ether)[59,65], poly (2-isopropyl-2-oxazoline), poly ((2-dimethylamino) ethyl methacrylate), poly (propylene oxide)[65] and poly (N-acryloylpiperidine)[56]. There are also polymers derived from natural sources with LCST behavior, such as methylcellulose, ethyl(hydroxyethyl)cellulose and hydroxypropylcellulose, that are of great interest for biomedical applications[65].

Some of thermoresponsive polymer structures are presented in Figure 1.

Despite the vast variety of thermoresponsive polymers, poly (N-isopropylacrylamide) (PNIPAM) is the most extensively studied one, as seen by the huge amount of publications dealing with this polymer. PNIPAM presents a LCST that lies between 30 and 35°C, depending on the precise microstructure of the macromolecule[77]. PNIPAM brings together an abrupt and reversible thermosensitive nature[78], with biocompatibility and LCST close to the human body[24], which turns it especially attractive in biomedical applications[79].

At temperatures below the LCST of PNIPAM, water‑polymer interactions are dominant, leading to dissolution or swelling in water[78]. With rising temperature, the polymer-solvent hydrogen bonds are being disrupted, whereas polymer-polymer interactions are greatly increased, resulting in collapsed structures at the LCST[80,81]. Above the transition temperature, the globules are aggregated into a few PNIPAM chains at very dilute solutions, while higher 238

polymer concentrations result in colloidal dispersions or even macroscopic precipitates[23,82,83].

Furthermore, as a consequence of LCST dependence on the balance between attractive polymer-polymer and polymer-solvent interactions[24,65], the transition temperature of PNIPAM can be adjusted by changing its molecular weight and concentration, by means of copolymerization of NIPAM with other comonomers, as well as through and addition of salts, surfactants and co-solvents[24,62,65,82-92].

3. pH-sensitive Polymers pH-sensitive polymers experience abrupt alterations in their polymer-polymer and polymer-solvent interactions in response to small variations in the environmental pH[61].

This behavior is attributed to the presence of pendant weak basic or acid groups in the polymeric chains, that either accept or release protons, respectively, as a result of slight changes in the pH of the medium[62]. This occurs because the degree of ionization of weak acids or bases is highly modified by changing the pH around their pKa value. Then, a large alteration in the hydrodynamic volume of the polymeric chains takes place, as a result of the quick variation in the charges of the pendant groups[2,63].

There are many polymers responsive to the environmental pH (Figure 2), including synthetic ones, such as poly (acrylic acid), poly (2-ethyl acrylic acid), poly (N,N-diethyl aminoethyl methacrylate) and poly (vinyl imidazole)[2], and those from natural sources, such as alginate[88,89], carboxymethylcellulose[90,91] and chitosan[92-97]. The pH‑responsive polymers can also be classified as: (i) polyacids, which contain pendant weak acid groups, such as -COOH and -SO3H; (ii) polybases, that possess pendant weak basic groups (-NH2) in their chains[2,10]; or (iii) polyamphoterics, that bear both weak acid and basic groups[62].

When applying pH-responsive polymers, the polysaccharides are preferred in many applications, because they join pH‑sensitiveness with inherent biological properties. One of the most outstanding pH-sensitive biopolymer is chitosan[98,99], a polysaccharide which naturally occurs in certain fungi[100], but is extensively obtained by the deacetylation of chitin[58,101,102], which is extracted from the shells of crustaceans, from the exoskeleton of many arthropods and from some fungi[103].

Despite the massive annual production and easy availability, due to its poor solubility in almost all common solvents, chitin does not find practical applications, except for being a source for obtaining chitosan[102-105].

Chitosan is a linear copolymer composed of two repeating units i.e. N-acetyl-2-amino-2-D-glucopyranose and 2-amino‑2deoxy-D-glucopyranose, linked by β-(1→4)‑glycosidic bonds. Generally, when the content of 2-amino-2-deoxyD-glucopyranose (degree of deacetylation – DD) in the polysaccharide chain is higher than 50 %, it becomes soluble in an aqueous acidic medium, as a result of the protonation of its amino groups and, in these conditions, it is named chitosan[106,107]. As a polysaccharide, chitosan exhibits attractive properties such as biocompatibility and biodegradability. Also, its degradation products are non-toxic, non-immunogenic and non-carcinogenic[62,106].

Polímeros , 25(3), 237-246, 2015

Development of dual-sensitive smart polymers by grafting chitosan with poly (N-isopropylacrylamide): an overview

Figure 1. Chemical structure of some thermosensitive polymers: (a) poly (N-isopropylacrylamide); (b) poly (N,N-diethylacrylamide); (c) poly (N-ethylacrylamide); (d) poly (2-carboxyisopropylacrylamide); (e) poly (N-(L)-(1-hydroxymethyl) propylmethacrylamide; (f) poly (N-acryloyl-N’-propilpiperazine); (g) poly (N-vinylcaprolactam); (h) poly ((2-dimethylamino)ethyl methacrylate); (i) poly (2-ethyl-2-oxazoline); (j) poly (2-isopropyl-2-oxazoline); (k) poly (vinyl methyl ether); (l) poly (propylene oxide); (m) methylcellulose; (n) ethyl(hydroxylethyl)cellulose.

The pH-responsive property of chitosan is a consequence of protonation–deprotonation equilibrium of its amino groups (pKa around 6) in aqueous media[58,102,108]. Besides, the presence of -OH and -NH2 reactive groups makes of this polysaccharide very attractive for chemical modifications[109,110].

4. Combination of pH and Thermosensitive Properties Through Graft Copolymerization of Chitosan with PNIPAM Nowadays, grafting PNIPAM side chains on the chitosan backbone constitutes a crescent area of research, since it bonds the most studied thermosensitive polymer with the most outstanding cationic polysaccharide, to reach dual Polímeros, 25(3), 237-246, 2015

temperature and pH responsive materials with remarkable properties[33,111-114].

To achieve those dual-responsive copolymers, researchers have applied some strategies. One of them involves a coupling reaction between PNIPAM bearing a reactive end group and chitosan, by using a condensing agent[115]. These condensing agents can catalyze the formation of amide bonds between carboxylic acid group of a carboxyl‑terminated PNIPAM and amine groups of chitosan. In these cases, an end‑functionalized PNIPAM has to be prepared before the graft copolymerization[116,117]. For instance, Rejinold et al. (2011) prepared PNIPAAM-COOH by using azobisisobutyronitrile (AIBN) and 3-mercaptopropionic acid, in isopropyl alcohol, at 75 °C; then, they grafted 239

Marques, N. N., Maia, A. M. S., & Balaban, R. C.

Figure 2. Chemical structure of some pH-responsive polymers: (a) poly (acrylic acid); (b) poly (methacrylic acid); (c) poly (2-ethyl acrylic acid); (d) poly (N,N-dimethyl aminoethyl methacrylate); (e) poly (N,N-diethyl aminoethyl methacrylate); (f) poly (vinyl imidazole); (g) chitosan; (h) alginate.

PNIPAAM-COOH chains onto chitosan backbone by using 1-ethyl-3-(3-dimethylaminopropyl) carbodiimide hydrochloride/N-hydroxysuccinimide (EDC/NHS) as the condensing agents, in acid medium, at room temperature.

The resulting nanoparticles were loaded with curcumin.

The in vitro drug release was effective only above LCST, which was attributed to higher polymer-polymer interaction than polymer-drug interaction when phase transition was reached. The drug loaded nanoparticles also showed cell uptake, cytocompatibility and specific toxicity on cancer cells, indicating that theses sensitive materials could be effective nanovehicles for controlled curcumin delivery[118].

Some others research groups have been preparing poly (N-isopropylacrylamide)-co-poly (acrylic acid) (PNIPAM‑co‑PAA) copolymers to be further anchored onto chitosan backbone by coupling reactions[119,120]. Proceeding in this method, PNIPAM-co-PAA copolymers were prepared 240

by the redox pair ammonium persulfate/ N,N,N’,N’tetramethylethylenediamine (APS/TEMED). Then, the PNIPAM-co-PAA side chains were bonded to chitosan by a coupling reaction, with the aid of EDC. The final copolymer was evaluated by means of its ability in removing phenol of aqueous solution. Phenol was oxidized by enzymatic reaction, aiming to produce compounds that could react with the amino groups of the chitosan derivative, forming Schiff bases or Michael-type adducts. Through heating and shaking the solution, the copolymer containing highly concentrated oxidized compounds deposited and agglutinated to a condensed coagulate. By increasing polymer concentration and chitosan content in the copolymer, the removal of phenol and its oxidized compounds was increased[119].

Core-shell nanoparticles, based on PNIPAM-co-PAA as core and chitosan as shell were designed, by means of previously copolymerization of N-isopropylacrylamide (NIPAM) with Polímeros , 25(3), 237-246, 2015

Development of dual-sensitive smart polymers by grafting chitosan with poly (N-isopropylacrylamide): an overview acrylic acid (AA), using sulfate persulfate as initiator and N,N-methylenebisacrylamide (MBA) as crosslink agent, at 75 °C in water medium, followed by coupling reaction between PNIPAM-co-PAA with chitosan, using EDC as the condensing agent, at 25 °C in aqueous media. The particles size was reduced from 380 to 25 nm as the temperature of the medium increased. While PNIPAM‑co‑PAA did not present thermosensitivity, core-shell smart nanoparticles showed temperature responsiveness and might also be more biocompatible than PNIPAM-co-PAA itself due to polysaccharide shell[121].

Chitosan-g-PNIPAM smart copolymers have also been prepared through radiation-based methods[80,122-124]. Zhao and collaborators prepared pH and temperature sensitive smart hydrogels by exposing the mixture of allylated chitosan, NIPAM and the photoinitiator, 2,2-dimethoxy2-phenylacetophenone (DMPA), in acid medium, to UV irradiation[125]. Swelling kinetics was dependent on pH, temperature and composition of the hydrogels. The in vitro release of the model drug methyl orange (MO) from the hydrogels was strongly pH dependent, being gradually released at pH 7.4 and rather low released in pH 2.0. This occurred due to the lack of ionic attractive interactions between MO and hydrogels at pH 7.4 and strong ionic attractive interactions between –SO3− groups of MO molecules and –NH3 + groups of chitosan at acid medium[125].

Grafting vinyl monomers onto the polysaccharide backbone using free radical polymerization (FRP) is a very common route to obtain chitosan-g-poly (N-isopropylacrylamide) responsive copolymers. Ceric ammonium nitrate (CAN), 2,2’-azoisobutyronitrile (AIBN) and persulfates (XPS) are some of the most employed initiators, which can be thermally activated, or using a redox initiation system[58,115].

There are many proposals regarding the anchored points for grafted chains when free radicals are used as initiators. When using CAN, for instance, some authors have exhibited a mechanism in which the chitosan units are predominantly oxidized through C2–C3 bond cleavage induced by Ce+4 ions, producing free-radicals sites onto the polysaccharide[126,127]. Others researchers suggest that the graft copolymerization of chitosan in the presence of CAN occurs onto amino groups of chitosan[78,128,129]. Initiation by persulfate has been also presented as occurring at different sites of chitosan backbone[130-133].

Duan and collaborators prepared chitosan-g-PNIPAM nanogels via free radical copolymerization at 80 °C, using APS as initiator and MBA as a crosslink agent. They suggested a synthetic route in which sulfate anion radicals, produced by thermal homolytic cleavage of APS, interacted with the hydroxyl groups of the polysaccharide to form alkoxy radicals, which then initiated the graft copolymerization of N-isopropylacrylamide (NIPAM) onto the backbone with MBA as a crosslinking agent. The final nanogels were loaded with oridonin (ORI), a powerful anticancer agent in chinese traditional medicine. The in vitro tests, performed at 37 °C, demonstrated a much faster drug release at acid condition than in pH 7.4. ORI loaded nanogels also presented better anti-tumor activity under acid media, as showed by both MTT assay and cellular morphological analysis, indicating Polímeros, 25(3), 237-246, 2015

that these nanogels are good candidates for pH-sensitive drug release of hydrophobic anticancer drugs, such as ORI[134].

More recently, our group evaluated the stability and rheological behavior of suspensions of PNIPAM, chitosan‑g‑PNIPAM and chitosan-g-(PNIPAM-co-PAA) particles, which were prepared by using potassium persulfate (50 °C) as initiator and MBA as a crosslinking agent.

Differences on particle-particle and particle-solvent attractive interactions were obtained by changing the composition of the particles and also pH and temperature environment, demonstrating that the particles stability can be adjusted depending on the desired application. The presence of chitosan onto the chemical network had particular importance on the particles behavior, as even at high pHs, in which chitosan is not protonated, the rigidity of polysaccharide chains helped to control stability of the particles[135].

Great attention has also been paid on controlled/living radical polymerization methods (CLRP), which include, mainly, atom transfer radical polymerization (ATRP), reversible addition fragmentation transfer (RAFT) and stable free radical polymerization (SFRP). The CLRP methods are based on a dynamic equilibrium between active species and dormant species, aiming to minimize the chance of termination reactions during the polymerization by decreasing the concentration of active species and, as a result, being able to produce polymers with precise architectures and compositions[58,136-138].

Chen and collaborators synthesized dual pH and temperature responsive chitosan-g-PNIPAM copolymers via ATRP. In order to protect the amino groups of chitosan, N-phthaloyl chitosan (PHCS) was firstly prepared, followed by the synthesis of the macroinitiator bromoisobutyryl‑terminated N-phthaloyl chitosan (PHCS-Br), through the reaction of PHCS with 2-bromoisobutyryl bromide, in the presence of triethylamine, in dimethylfomamide (DMF) medium; then, PHCS-g-PNIPAM was prepared by combining PHCS-Br macroinitiatior with NIPAM, 2,2’-Bipyridyl and cuprous chloride (CuCl), in DMF at 70 °C and under nitrogen atmosphere; finally, deprotection was made by the reaction of PHCS-g-PNIPAM with hydrazine hydrate in water medium, under nitrogen atmosphere. The LCST of chitosang-PNIPAM in aqueous solution was 33 °C at pH 6.3 and 35 °C at pH 5.0, which indicated that the thermoresponsive behavior is also pH dependent in these materials, since the lowest pH implies on a more hydrophilic material[139].

5. Conclusions In this paper, the stimuli-sensitive polymer systems are described, by means of their behavior and applications.

The pH-responsive, thermoresponsive and the dual pH and thermoresponsive copolymers were presented, giving focus on the combination of chitosan with poly (N-isopropylacrylamide) by graft copolymerization, obtained via different routes and their properties. Chitosan-g-PNIPAM copolymers are very promising materials especially on biomedical applications, such as in drug delivery systems and tissue engineering.

The pH-sensitiveness and biological favorable properties from chitosan associated to the thermosensitive properties from PNIPAM lead to a powerful responsive material that 241

Marques, N. N., Maia, A. M. S., & Balaban, R. C. can be synthesized to a target physical-chemical behavior, as well as with an enhanced therapeutic efficiency and reduced side effects.

6. Acknowledgements The authors are grateful to CAPES, a Brazilian Government entity targeting the training of human resources, for their financial support.

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135. Nascimento, N. N., Curti, P. S., Maia, A. M. S., & Balaban, R. C. (2013). Temperature and pH effects on the stability and rheological behavior of the aqueous suspensions of smart polymers based on N-isopropylacrylamide, chitosan, and acrylic acid. Journal of Applied Polymer Science, 129(1), 334-345. http://dx.doi.org/10.1002/app.38750.

136. Oh, J. K., Lee, D. I., & Park, J. M. (2009). Biopolymer-based microgels/nanogels for drug delivery applications. Progress in Polymer Science, 34(12), 1261-1282. http://dx.doi.org/10.1016/j. progpolymsci.2009.08.001.

137. Cayre, O. J., Chagneux, N., & Biggs, S. (2011). Stimulus responsive core-shell nanoparticles: Synthesis and applications of polymer based aqueous systems. Soft Matter, 7(6), 22112234. http://dx.doi.org/10.1039/C0SM01072C. 138. Braunecker, W. A., & Matyjaszewski, K. (2007). Controlled/ living radical polymerization: Features, developments, and perspectives. Progress in Polymer Science, 32(1), 93-146. http://dx.doi.org/10.1016/j.progpolymsci.2006.11.002.

139. Chen, C., Liu, M., Gao, C., Lü, S., Chen, J., Yu, X., Ding, E., Yu, C., Guo, J., & Cui, G. (2013). A convenient way to synthesize comb-shaped chitosan-graft-poly (N-isopropylacrylamide) copolymer. Carbohydrate Polymers, 92(1), 621-628. http:// dx.doi.org/10.1016/j.carbpol.2012.09.014. PMid:23218344.

Received: Apr. 10, 2014 Revised: Aug. 19, 2014 Accepted: Nov. 14, 2014 Polímeros , 25(3), 237-246, 2015

http://dx.doi.org/10.1590/0104-1428.1819

TG/FT-IR characterization of additives typically employed in EPDM formulations Natália Beck Sanches1, Silvana Navarro Cassu1,2 and Rita de Cássia Lazzarini Dutra1,2 Instituto Tecnológico de Aeronáutica – ITA, CEP 12228-900, São José dos Campos, SP, Brazil 2 Divisão de Química – AQI, Instituto de Aeronáutica e Espaço – IAE, CEP 12228-904, São José dos Campos, SP, Brazil

1

*ritalazzarini@yahoo.com.br

Abstract Thermogravimetric analysis coupled to Fourier transform infrared spectroscopy (TG/FT-IR) is a very popular technique for rubbers characterization. It involves analyses of the base polymer and additives. Ethylene–propylene–diene (EPDM) rubbers are frequently investigated by TG/FT-IR; however, the focus has been the degradation temperature range of the polymer. In this study, unvulcanized and vulcanized EPDM rubber and its additives were investigated

by TG/FT-IR, without solvent extraction, and in a wide temperature range. Initially, the additives were individually characterized. TG/FT-IR identified the characteristic groups of all the additives analyzed and distinguished them from each other. Afterwards, unvulcanized and vulcanized EPDM rubbers were investigated without prior extraction.

TG/FT-IR detected absorptions due to the additives tetramethylthiuram monosulfide and 2-mercaptobenzothiazole. Both of these sulfur‑containing additives were present in the EPDM formulation at concentrations of 0.7 phr (0.63 wt %).

The TG/FT-IR technique had some limitations, because not all the additives in EPDM rubber were detected. Paraffin oil, stearic acid and 2,2,4-trimethyl-1,2-dihydroquinoline functional groups were not observed in either the unvulcanized or vulcanized EPDM. Nevertheless, in addition to the ability of this method to detect sulfur-containing groups, the lack of a pre-extraction reduces the time and effort required for additive analysis in rubbers.

Keywords: EPDM, additives, characterization, TG/FT-IR, TMTM, MBT.

1. Introduction Additives are selected and incorporated into rubbers to provide specific properties. Useful rubbers can only be obtained by appropriate compounding. Some chemicals provide processing aid, extended shelf-life or improved long‑term performance, others enhance polymer properties.

As a result, rubbers are complex chemical materials, which are difficult to analyze[1]. Analytical techniques that enable the detection of additives are of great importance for industries, especially for those cases in which they are present in very low concentrations.

Typically, there are two approaches for additive analysis in rubbers: extraction with solvent prior to analysis or direct determination. Extraction procedures can be very complex, labor-intensive, and not always reproducible[1-3]; thus, direct analysis is always preferred if it is feasible.

Among the techniques for direct determination, thermogravimetric analysis coupled with Fourier transform infrared spectroscopy (TG/FT-IR) is one of the most powerful methods to study the thermal degradation of polymers.

This technique has the merit of identifying the evolved gases of a polymer at different degradation temperatures; thus, it allows a temperature selective analysis[4].

Ethylene–propylene–diene copolymer (EPDM) is one of the most important rubbers, and has uses in diverse applications, even in the aerospace field. Its saturated backbone provides remarkable resistance to oxygen, ozone, and heat[5].

Some recently published studies have been related to the TG/FT‑IR analysis of additives in EPDM. Jiang et al.[6] used

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TG/FT-IR to evaluate the effect of polyphenylsilsesquioxane (PPSQ) on the release of volatile products in EPDM samples.

The results indicated that PPSQ affects volatile products of EPDM and is detected in its formulation by TG/FT-IR. Çavdar et al.[7] studied different vulcanizing agent contents

by TG/FT-IR. They observed that increasing vulcanizing agent content decreased band intensities of CO and CO2 and enhanced the thermal stability of EPDM rubber. Özdemir[8] utilized TG/FT-IR to evaluate irradiated EPDM rubber vulcanized with two types of peroxides. The main absorptions of this irradiated rubber were attributed to aromatic C-H, methylene C-H, methyl ether C-H, methyl C-H, CO, and CO2. However, none of the studies attempted to detect additives at degradation temperatures other than the EPDM polymer degradation temperature.

In a previous study[9], we employed Fourier transform infrared spectroscopy of gaseous pyrolyzates (PY-G/FT-IR) for the detection of additives in EPDM rubber. The absorptions of additives were identified in unvulcanized and vulcanized EPDM samples without prior extraction with solvent.

This technique was able to detect sulfur-containing additives at concentrations as low as 1.4 phr (1.26%). However, as the whole amount of the evolved gas from pyrolysis was trapped into a gas cell at once, a temperature selective analysis was not possible.

A temperature selective analysis can be performed using TG/FT-IR; moreover, this technique can cover a wide range of temperatures and provide information related

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S S S S S S S S S S S S S S S S S S S S

Sanches, N. B., Cassu, S. N., & Dutra, R. C. L. to evolving products. Hence, this study is aimed at the detection of additives in EPDM rubber using TG/FT-IR without solvent extraction. Initially, each additive was individually characterized to identify its characteristic absorptions. Subsequently, unvulcanized and vulcanized EPDM samples were evaluated using raw EPDM rubber as a reference sample. The absorptions related to additives were identified by comparing the TG/FT-IR spectra of additives and EPDM samples.

2. Materials and Methods 2.1. Materials Paraffin oil, stearic acid, 2,2,4-trimethyl-1,2-dihydroquinoline (TMQ), tetramethylthiuram monosulfide (TMTM), and 2-mercaptobenzothiazole (MBT) were provided by Zanaflex Borrachas Ltda, Brazil, and analyzed individually as received.

As a reference sample, commercial grade EPDM Keltan 21 containing the diene ENB (ethylidene norbornene) was purchased from DSM Elastômeros do Brasil Ltda, Brazil, and used as received.

Unvulcanized and vulcanized (terminology in accordance with the ASTM D1566-11[10]) samples were prepared using EPDM Keltan 21 and rubber-grade chemicals according to the composition listed in Table 1.

The molecular structures of the analyzed compounds prior to degradation are shown in Figure 1. MBT is represented by two different structures because this additive can be present in two tautomeric forms (Figures 1v and 1vi)[11-13].

Contini et al.[12] stated that MBT exists in the vapor phase only in its tautomeric thione (benzothiazoline-2-thione) form, which contains a C=S bond and a hydrogen bonded to the nitrogen, rather than the thiol form, which contains an endocyclic C=N bond and a hydrogen bonded to sulfur.

According to Wu et al.[13] and Mohamed et al.[14], thione is the dominant form in the solid state. For a better understanding of the structural differences, the thione form of the MBT molecule is shown in Figure 1.

2.2. TG/FT-IR analysis TG/FT-IR analyses were performed using a PerkinElmer Pyris 1 TGA coupled with a PerkinElmer Spectrum One FT‑IR. The transfer line and FT-IR gas cell were maintained at 210 and 230 °C, respectively. The spectra were collected at resolution of 8 cm−1 with a co-add of 8 scans per spectrum, resulting in one spectrum collected every 22 s. Each sample (ca. 20 mg) was heated from room temperature to 900 °C under a nitrogen atmosphere.

Heating rate and gas flow are the experimental parameters that have the greatest effect on the results of a TG/FT-IR experiment[15]. Based on previous studies and recommendations of the manufacturer of the equipment, tests were performed to determine the optimal settings.

Moreover, Gram–Schmidt reconstruction (GSR) profiles, which indicate the relative intensities of gases in the gas cell, were qualitatively compared. The optimal conditions, which maximized the FT-IR response, were determined to be the combination of a 20 °C/min heating rate with a 25 ml/min gas flow. According to Berbenni et al.[15], a 248

well-adjusted TG/FT-IR experiment will present GSR and differential thermogravimetric (DTG) curves with similar profiles. In addition to providing the most intense absorptions, the selected parameters also showed excellent agreement between the GSR and DTG curves.

2.3. Selection of FT-IR spectra In a typical TG/DTG result, the maximum value of the derivative curve indicates the maximum degradation rate.

At this value, the quantity of evolving gases reaches its maximum. A single TG/DTG experiment can present one or multiple maxima depending on the sample components and its degradation characteristics.

An FT-IR spectrum was collected for each observed maximum degradation rate, which agreed with the maxima of the GSR profile with a delay of a few seconds. For some samples, a single spectrum was sufficient to represent the whole experiment, whereas for other samples, in which more than one degradation event was detected, more spectra were collected.

Moreover, the main absorptions determined in the spectra of additives were investigated in unvulcanized and vulcanized EPDM samples at the same temperature at which they appeared in the degradation of the additives.

Table 1. Composition of EPDM rubber (partially reproduced from Sanches et al.[9]).

Component EPDM Keltan 21 Paraffin oil Stearic acid TMQ (2,2,4-trimethyl-1,2-dihydroquinoline) TMTM (tetramethylthiuram monosulfide) MBT (2-mercaptobenzothiazole) ZnO (zinc oxide) Carbon black S (sulfur) a

Unvulcanized and Vulcanized (phra) 100 1.0 0.5 1.0 0.7 0.7 2.0 5.0 0.7

parts per hundred parts of rubber.

Figure 1. Molecular structure of the analyzed compounds prior to degradation (adapted from Sanches et al.[9]).

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TG/FT-IR characterization of additives typically employed in EPDM formulations In this study, all the FT-IR spectra were obtained from TG/FT-IR experiments, which are referred to as TG/FT-IR spectra.

3. Results and Discussion 3.1. Degradation temperatures Initially, the additives, raw EPDM, and both unvulcanized and vulcanized EPDM were analyzed separately by TG/FT-IR. A set of TG/DTG curves, GSR curves, and FT-IR spectra was obtained for each sample.

Figure 2 shows the degradation temperatures obtained from the TG/DTG curves. A comparison between the degradation ranges of EPDM and its additives is useful because it indicates the temperature at which the characteristic absorptions of the additives should be searched for in the unvulcanized and vulcanized EPDM TG/FT-IR experiments.

Figure 2 shows that only TMTM presents a narrow degradation temperature region of 180-310 °C, whereas the other additives show wide degradation regions. Raw EPDM degradation starts at approximately 250 °C; therefore, at lower temperatures, additive absorptions should be observed in unvulcanized and vulcanized EPDM without interference from characteristic EPDM bands. At higher temperatures, most additive absorptions should simultaneously be observed with the polymer bands.

The comparison between raw, unvulcanized and vulcanized EPDM confirm their peculiarities. Raw EPDM presents a more narrow temperature degradation range, as expected for a neat polymer. Unvulcanized and vulcanized EPDM degradation is broader because of the additive content. The comparison between unvulcanized and vulcanized EPDM shows that unvulcanized EPDM degradation begins at a lower temperature, although both contain the same formulation. Vulcanized EPDM was heated

to approximately 150-180 °C in the crosslinking process; therefore, to some extent, it loses a certain quantity of low molecular weight additives.

3.2. TG/FT-IR analysis of additives IR spectra of gaseous products can be very complex because they present a large number of absorptions. In this study, for peak assignment, the presence or absence of characteristic functional groups in the TG/FT-IR spectra was used. Figure 3 shows the FT-IR spectra of additives obtained from the TG/FT-IR technique.

The main bands observed in the TG/FT-IR spectrum of paraffin oil (Figure 3a) are 3085, 3016, 2933, 2868, 1462, 1380, 949, and 911 cm−1. The peaks at 3085 and 3016 cm−1 are assigned to the C-H group and/or C-H aromatic group, whereas those at 2933, 2868 and 1380 cm−1 are assigned to the CH3 group. The band at 1462 cm−1 is assigned to C-H group. Although aromatic groups are not expected in the TG/FT-IR spectrum of paraffin oil, its presence can be explained by the fact that rubber-grade paraffin oils may contain 26%-40% of naphthenic oil and 2%-7% of aromatic oil[16].

The TG/FT-IR spectrum of stearic acid is shown in Figure 3b. The band at 3576 cm−1 is assigned to the OH group. The peaks at 2933 and 2864 cm−1 are assigned to the CH2 groups. The absorption at 1776 cm−1 is assigned to the C=O group. The bands at 1462 and 1126 cm−1 are assigned to the CH2 group, whereas the band at 1372 cm−1 is assigned to the CH3 group[17,18].

The TG/FT-IR spectra corresponding to TMQ degradation at 362 and 473 °C are shown in Figures 3c and 3d, respectively.

The spectrum at 362 °C showed a band at 3015 cm−1, which is assigned to the C-H or H-C=C group. The bands around 2968–2877 and 1059-1005 cm−1 are assigned to the CH3 group. The peaks at 1603 and 744 cm−1 are assigned to the

Figure 2. Degradation temperature of components from TG/FT-IR experiments.

Polímeros, 25(3), 247-255, 2015

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Sanches, N. B., Cassu, S. N., & Dutra, R. C. L.

Figure 3. TG/FT-IR spectra of (a) paraffin oil at 450 °C; (b) stearic acid at 324 °C; (c) TMQ at 362 °C; and (d) TMQ at 473 °C.

C-C aromatic and C-H groups, respectively. The absorption at 1304 cm−1 is assigned to the N-H[19] and/or C-N[20] groups, whereas the bands at 1263 and 1162 cm−1 are assigned to the C-N aromatic group.

The spectrum at 473 °C showed the same bands as the one at 362 °C; however, other absorptions were detected, which are described as follows. The bands at 1497, 1380, and 835-812 cm−1 are assigned to the C-C aromatic and/or CHN, CH3, and C-H groups of the benzene ring, respectively[20].

These bands were possibly not detected at 362 °C because of the low amount of evolved gas at this temperature.

Figure 4a shows the TG/FT-IR spectrum of TMTM at 294 °C. The bands at 2941-2812 cm−1 are assigned to the CH3 group. The bands at 2072 and 2048 cm−1 probably can be assigned to the N=C=S (isothiocyanate) group[21-24].

The spectra of CS2 in the gaseous state obtained from the reference databases[25,26], show absorptions at approximately 2320, 2179, and 1530 cm−1, which are in excellent agreement with the bands observed at 2336-2316, 2193-2179, and 1539-1524 cm−1. Moreover, similar to other studies, a band located between 1523 and 1541 cm−1 can be attributed to the CS2 group[27,28].

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Therefore, considering the chemical structure of TMTM and the references in the literature, doublets can be attributed to the presence of the C=S and/or CS2 group in the TMTM degradation products.

The TG/FT-IR spectra corresponding to MBT degradation at 361 and 784 °C are shown in Figures 4b and 4c, respectively.

The spectrum at 361 °C has a band at 3748 cm−1, which can be assigned to the N-H group. The peaks at 3076 and 755 cm−1 are assigned to the aromatic C-H group. The band at 2895 cm−1 is assigned to the CH2 group. The doublet at 2071-2046 cm−1 and the peak at 656 cm−1 are assigned to the N=C=S group[20]. These absorptions are in excellent agreement with the benzothiazole spectrum in the literature[25].

Doublets at 2193-2179 (very subtle) and 1539-1525 cm−1 are assigned to the C=S and/or CS2 group, which is analogous to the TMTM assignment. These assignments confirm the presence of the thione form in the vapor phase, as indicated

by Contini et al.[12].

After 450 °C, the intensity of the bands observed in the 361 °C spectrum starts decreasing with the emergence of a doublet at 1376-1343 cm−1, indicating the structural transformation of the molecule. The spectrum at 784 °C shows this doublet; however, it is inconclusive for determining the evolved products during MBT thermal degradation.

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TG/FT-IR characterization of additives typically employed in EPDM formulations

Figure 4. TG/FT-IR of (a) TMTM at 294 °C; (b) MBT at 361 °C; and (c) MBT at 784 °C.

According to Brooks et al.[29], pyrolysis of pure benzene leads to ring opening at 763 °C, leading to the generation of methane as one of the degradation products. The spectrum of methane from the literature[25] shows absorptions in the region of 1376-1343 cm−1, indicating that they could be related to the CH3 group. Nevertheless, recent studies have assigned these peaks to the C-C group of the benzene ring[30,31], or to the ring vibrations of the heterocyclic MBT ring[21].

3.3. TG/FT-IR analysis of EPDM TG/FT-IR results of raw, unvulcanized, and vulcanized EPDM can complement each other. In this study, raw rubber was analyzed as a reference sample, and its FT-IR spectrum was compared with the spectra of unvulcanized and vulcanized EPDM to differentiate the polymer absorptions.

Unvulcanized rubber was analyzed by TG/FT-IR to obtain the spectrum before vulcanization, which is when additives are chemically preserved.

The TG/FT-IR spectra of the evolved products of raw, unvulcanized, and vulcanized EPDM are shown in Figure 5.

Figure 5a shows the TG/FT-IR spectrum of raw EPDM. The peak at 3086 cm−1 is assigned to the olefinic and/or aromatic C-H. The peaks at 988 and 911 cm−1 are assigned to the vinylic C=C. The band around 949 cm−1 is assigned Polímeros, 25(3), 247-255, 2015

to the trans C=C. The bands observed at 889 cm−1 and at 1385 cm−1 are assigned to the RR′CCH2 and CH3 groups, respectively[20]. The absorptions detected in the raw EPDM FT-IR spectrum can help evaluate additive-related absorptions in the formulated EPDM. TG/FT-IR spectra of unvulcanized and vulcanized EPDM (Figures 5b and 5c) are very similar at 272 and 280 °C, respectively. These spectra show very weak doublets around 2192-2179 and 2071-2047 cm−1. A more intense doublet is detected in the region of 1540-1525 cm−1. These peaks are absent in the spectrum of raw EPDM (Figure 5a); thus, they are related to the additives. Only TMTM and MBT TG/FT-IR spectra show similar bands; thus, it can be assumed that these absorptions are related to them.

Figures 5d and 5e show the TG/FT-IR spectra of unvulcanized and vulcanized EPDM at 509 and 517 °C, respectively. At these temperatures, unvulcanized and vulcanized EPDM degradation products show only polymerrelated absorptions, which are in excellent agreement with the raw EPDM TG/FT-IR spectrum (Figure 5a).

TG/FT-IR results of raw EPDM, as well as the results for unvulcanized and vulcanized EPDM at 509 and 517 °C, respectively, exclusively show the EPDM polymer.

Their assignments are in accordance with a published study[32] 251

Sanches, N. B., Cassu, S. N., & Dutra, R. C. L.

Figure 5. TG/FT-IR of (a) raw EPDM at 494 °C; (b) unvulcanized EPDM at 272 °C; (c) vulcanized EPDM at 280 °C; (d) unvulcanized EPDM at 509 °C; and (e) vulcanized.

that evaluated EPDM rubber by coupling pyrolysis–gas chromatography and mass spectrometry (PY-GC/MS).

Its results indicate a mix of alkanes and alkenes among the major products of EPDM thermal degradation.

Moreover, the main absorptions of paraffin oil, stearic acid, and TMQ were investigated in EPDM TG/FT-IR spectra at the temperature range, in which the additive degradation 252

was observed. The TG/FT-IR spectra of unvulcanized and vulcanized EPDM show no peaks that could be related to these additives.

In a previous study[9], a band at 771/772 cm−1 from paraffin oil, stearic acid, and TMQ pyrolysis was detected in the PY-G/FT-IR spectra of unvulcanized and vulcanized EPDM. In the TG/FT-IR experiments, the absence can be Polímeros , 25(3), 247-255, 2015

TG/FT-IR characterization of additives typically employed in EPDM formulations explained by features of the technique. In this study, the evolved gases are carried to FT-IR by the flow gas, which dilutes the degradation products; moreover, the amount of sample is significantly less than the one analyzed with the PY-G/FT-IR technique.

Table 2 summarizes the functional groups assigned to the additives and EPDM by TG/FT-IR experiments, and relates them to the temperature at which spectra were collected. It can be seen that unvulcanized and vulcanized EPDM spectra were collected at a higher temperature than the raw EPDM spectrum. As the temperatures are related to the maximum degradation rate of the polymer, the differences are possibly

due to the additives (unvulcanized and vulcanized EPDM) and the crosslinking effect (vulcanized EPDM only).

Some differences in absorption intensities are expected between the TG/FT-IR spectra of unvulcanized and vulcanized EPDM. Vulcanized gaseous output is expected to be lower in quantity with less intense absorptions because the crosslinking process involves temperatures around 150 °C, which can alter the additives and cause chemical reactions.

Although this study involves only qualitative analysis, this quantitative aspect could prevent detection of additives.

For example, sulfur compounds could be partially or totally consumed in the reticulation reaction; consequently, their bands might lose intensity or even disappear from the

Table 2. Functional groups assigned in this study.

Component Paraffin Oil

Temperature 450 °C

Stearic Acid

324 °C

TMQ

362 °C

473 °C

TMTM

MBT

294 °C

361 °C

784 °C Raw EPDM

494 ºC

Unvulcanized and vulcanized EPDM

272 and 280 °C, respectively 509 and 517 °C, respectively

Polímeros, 25(3), 247-255, 2015

Wavenumber (cm–1) 2933, 2968 and 1380 3085 and 3016

Functional Groups assigned by TG/FT-IR CH3 CH and/or C-H aromatic

1462 3576 2933 and 2864 1776 1462 and 1126 1372 3015 2968-2877 and 1059-1005 1063 744 1304 1263 and 1162 3015 2968-2877 and 1059-1005 1063 744 1304 1263 and 1162 1380 835-812 2336-2316, 2193-2179 and 1539-1524 2941–2812 2072 and 2048 3748 3076 and 755 2895 2071-2046 and 656 2193-2179 and 1539-1525 1376-1343

CH OH CH2 C=O CH2 CH3 C-H or H-C=C CH3 C-C aromatic C-H N-H and/or C-N C-N aromatic C-H or H-C=C CH3 C-C aromatic C-H N-H and/or C-N C-N aromatic CH3 C-H of benzene ring CS2 and/or C=S

988 and 911 949 889 1385 3086 2192-2179,

CH3 N=C=S N-H C-H CH2 N=C=S CS2 and/or C=S CH3 or C-C aromatic or ring vibrations of the heterocyclic ring C=C vinyl C=C trans RR′CCH2 CH3 C-H olefinic and/or aromatic CS2 and/or C=S

2071-2047 and 1540-1525 Same wavenumbers as raw EPDM

Same absorptions as raw EPDM

253

Sanches, N. B., Cassu, S. N., & Dutra, R. C. L. FT‑IR spectra of vulcanized rubber. Nevertheless, in this study, the unvulcanized and vulcanized EPDM TG/FT-IR spectra are very similar.

Intriguingly, the TG/FT-IR technique was able to detect only the sulfur-containing additives TMTM and MBT in unvulcanized and vulcanized EPDM. Both additives demonstrated similar characteristic absorptions related to sulfur compounds in the regions of 2192-2179, 2071-2047, and 1540-1524 cm−1. These results confirm that absorptions related to the sulfur compounds show stronger intensities in the gaseous state than in the solid and liquid state[9].

TG/FT-IR was capable of detecting sulfur compounds in concentrations as low as 1.4 phr (1.26%), considering the sum of TMTM and MBT content in the compounds.

This confirms the potential of this method to investigate these types of materials, even with a considerably smaller sample than that used in the PY-G/FT-IR technique[9].

In our former study[9], the PY-G/FT-IR technique was unable to differentiate TMTM and MBT from each other; however, TG/FT-IR could easily distinguish between them.

This demonstrates the superiority of TG/FT-IR for analyzing sulfur additives. In fact, PY-G/FT-IR and TG/FT-IR can complement each other because the former has the advantage of providing a very concentrated evolved gas, which can be helpful in analyzing additives whose products present weak absorptions or are present in low concentrations in the rubber compound. On the other hand, TG/FT-IR has the disadvantage of yielding gases diluted with the flow gas, but is able to provide spectra at different temperatures and thus differentiate the components of complex degradation samples.

4. Conclusion TG/FT-IR characterization of additives frequently used in EPDM rubber was performed. This technique was capable of distinguishing all the additives from each other without exception. By analyzing the degradation of these additives at different temperatures, TG/FT-IR demonstrated very distinct spectra. Furthermore, TMTM and MBT accelerators showed some similar characteristic absorptions, which are related to sulfur compounds. The presence of these specific absorptions enables the differentiation of additives with and without sulfur, by TG/FT-IR. This technique was able to detect absorptions of sulfur additives in EPDM rubber at concentrations as low as 1.4 phr (1.26%), even in vulcanized EPDM. These specific absorptions were detected at temperatures lower than the temperatures at which the polymer bands were observed, in accordance to TMTM and MBT degradation characteristics.

Moreover, the identification of functional groups of these additives was possible without their prior extraction using solvents in both unvulcanized and vulcanized EPDM. Therefore, the TG/FT-IR technique can be employed for the analysis of separate additives, neat rubber (raw), and unvulcanized and vulcanized compounds. It demonstrated temperature selectivity, enabling the investigation of specific temperatures at which the additives degrade. Although the technique was unable to detect paraffin oil, stearic acid, and TMQ additives in unvulcanized and vulcanized EPDM, the sulfur-related characteristic absorptions of TMTM and MBT were identified.

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9. Sanches, N. B., Cassu, S. N., Diniz, M. F., & Dutra, R. C. L. (2014). Characterization of additives typically employed in EPDM formulations by using FT-IR of gaseous pyrolyzates.

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10. American Society for Testing and Materials (2011). ASTM D1566-11: Standard Terminology Relating to Rubber. West Conshohocken: ASTM. 11. Nieuwenhuizen, P. J., Reedijk, J., Van Duin, M., & McGill, W. J. (1997). Thiuram-and dithiocarbamate-accelerated sulfur vulcanization from the chemist’s perspective; methods, materials and mechanisms reviewed. Rubber Chemistry and Technology, 70(3), 368-429. http://dx.doi.org/10.5254/1.3538436.

12. Contini, G., Di Castro, V., Stranges, S., Richter, R., & Alagia, M. (2002). Gas-phase photoemission study of 2-Mercaptobenzothiazole. The Journal of Physical Chemistry A, 106(12), 2833-2837. http://dx.doi.org/10.1021/jp013423b.

13. Wu, F.-L., Hussein, W. M., Ross, B. P., & McGeary, R. P. (2012). 2-Mercaptobenzothiazole and its Derivatives: Syntheses, Reactions and Applications. Current Organic Chemistry, 16(13), 1555-1580. http://dx.doi.org/10.2174/138527212800840964.

14. Mohamed, T. A., Mustafa, A. M., Zoghaib, W. M., Afifi, M. S., Farag, R. S., & Badr, Y. (2008). Reinvestigation of benzothiazoline-2-thione and 2-mercaptobenzothiazole tautomers: conformational stability, barriers to internal rotation and DFT calculations. Journal of Molecular Structure: THEOCHEM, 868(1-3), 27-36. http://dx.doi.org/10.1016/j. theochem.2008.07.037.

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TG/FT-IR characterization of additives typically employed in EPDM formulations TG/Spectral response. Thermochimica Acta, 258, 125-133. http://dx.doi.org/10.1016/0040-6031(94)02237-I. 16. Loadman, M. J. R. (1998). Analysis of rubber and rubber-like polymers. Dordrecht: Klumer Academic Publishers.

17. Zhang, Z., Dallek, S., Vogt, R., Li, Y., Topping, T. D., Zhou, Y., Schoenung, J. M., & Lavernia, E. J. (2010). Degassing behavior of nanostructured al and its composites. Metallurgical and Materials Transactions. A, Physical Metallurgy and Materials Science, 41(2), 532-541. http://dx.doi.org/10.1007/ s11661-009-0089-6.

18. Jansen, J. A. J., Van Der Maas, J. H., & Boer, A. P. (1992).

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19. Welti, D., & Stephany, R. (1968). Some comments on the infrared spectra of vapors. Applied Spectroscopy, 22(6), 678688. http://dx.doi.org/10.1366/000370268774384119.

20. Smith, A. L. (1979). Applied infrared spectroscopy. New York: John Wiley & Sons.

21. Silverstein, R. M., Webster, F. X., & Kiemle, D. J. (2005).

Spectrometric identification of organic compounds. New York: John Wiley & Sons.

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23. Fiels, L. D., Sternhell, S., & Kalman, J. R. (2008). Organic structures from Spectra. Chichester: John Wiley & Sons.

24. Coates, J. (2000). Interpretation of infrared Spectra, a practical approach. In R. A. Meyers (Ed.), Encyclopedia of analytical chemistry (pp. 10815-10837). Chichester: John Wiley & Sons.

25. NIST Standard Reference Database WebBook. (2013). Retrieved in 3 Mar. 2013, from http://webbook.nist.gov/cgi/cbook. cgi?ID=C75150&Units=SI&Type=IR-SPEC

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26. BIO-RAD/SADTLER Reference Database. (2013). Retrieved in 3 Mar. 2013, from Retrieved from https://scifinder.cas.org/ 27. Person, W. B., & Hall, L. C. (1964). Absolute infrared intensities of CS2 fundamentals in gas and liquid phases: an interpretation of the bond moments of CO2 and CS2. Spectrochimica Acta, 20(5), 771-779. http://dx.doi.org/10.1016/0371-1951(64)80076-X. 28. Hoffmann, F., Riesen, R., & Foreman, J. (2000). Characterization of thermal stability and reaction products by means of TGAFTIR coupling. American Laboratory, 32(1), 13-17.

29. Brooks, C. T., Peacock, S. J., & Reuben, B. G. (1979). Pyrolysis of benzene. Journal of the Chemical Society, Faraday Transactions I, 75(0), 652-662. http://dx.doi.org/10.1039/f19797500652.

30. Xiao-Hong, L., Zheng-Xin, T., & Xian-Zhou, Z. (2009).

Molecular structure, IR spectra of 2- mercaptobenzothiazole and 2-mercaptobenzoxazole by density funcional theory and ab initio Hartree-Fock calculations. Spectrochimica Acta. Part A: Molecular and Biomolecular Spectroscopy, 74(1), 168-173. http://dx.doi.org/10.1016/j.saa.2009.05.026.

31. Sathyanarayanmoorthi, V., Karunathan, R., & Kannappan, V. (2013). Molecular modeling and spectroscopic studies of Benzothiazole. Journal of Chemistry, 2013, 1-14. http://dx.doi. org/10.1155/2013/258519.

32. Choi, S., & Kim, Y. (2011). Analysis of 5-ethylidene-2norbornene in ethylene-propylene-diene terpolymer using pyrolysis-GC/MS. Polymer Testing, 30(5), 509-514. http:// dx.doi.org/10.1016/j.polymertesting.2011.04.005.

Received: June 25, 2014 Revised: Oct. 08, 2014 Accepted: Nov. 17, 2014

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http://dx.doi.org/10.1590/0104-1428.1747

S S S S S S S S S S S S S S S S S S S S

Blends of ground tire rubber devulcanized by microwaves/ HDPE - Part A: influence of devulcanization process Fabiula Danielli Bastos de Sousa1, Júlia Rocha Gouveia1, Pedro Mario Franco de Camargo Filho1, Suel Eric Vidotti1, Carlos Henrique Scuracchio1, Leice Gonçalves Amurin2 and Ticiane Sanches Valera2 1 Centro de Engenharia, Modelagem e Ciências Sociais Aplicadas – CECS, Universidade Federal do ABC – UFABC, CEP 09210-580, Santo André, SP, Brazil 2 Departamento de Engenharia Metalúrgica e de Materiais, Escola Politécnica, Universidade de São Paulo – USP, CEP 05508-030, São Paulo, SP, Brazil

*fabiuladesousa@gmail.com

Abstract The main objective of this work is the study of the influence of microwaves devulcanization of the elastomeric phase on dynamically revulcanized blends based on Ground Tire Rubber (GTR)/High Density Polyethylene (HDPE).

The devulcanization of the GTR was performed in a system comprised of a conventional microwave oven adapted with a motorized stirring at a constant microwaves power and at various exposure times. The influence of the devulcanization process on the final properties of the blends was evaluated in terms of mechanical, viscoelastic, thermal and rheological properties. The morphology was also studied.

Keywords: elastomers, recycling, GTR, devulcanization, HDPE.

Introduction The search for new materials has been a constant in the human history. Similarly, solutions to the problem of disposal of waste polymers, especially waste rubber, that causes serious environmental problems and concern, have been desired for many years. Rubber requires a long period of time to degrade naturally due to its structure of cross‑linkings and the presence of stabilizers and other additives[1,2].

The technique of devulcanization by microwaves is currently one of the most promising ones, based on the good properties of the devulcanized material and the possibility of high productivity. The process takes advantage of volumetric heating of the material by microwaves, promoting a more uniform heating than that achieved by traditional methods of heating, which depend on conduction and/or convection[3-5].

Materials react differently when exposed to an electromagnetic field, like the one generated by microwaves. In dielectric materials, molecules or free ions are rearranged in dipole momentum which results in the volumetric heating through the volume of the material. These molecules vibrate at high frequency tending to re-orient and align themselves with the microwave field. Interaction between the material and the microwave energy generates heat. The ability to convert microwave energy into thermal energy depends on the magnitude of the dielectric loss of the material[6,7]. Therefore, in mixtures of materials, it is possible a selective heating of specific regions, a property that has been exploited in the processing of thermosets with mineral fillers[3-5]. Elastomers such as natural rubber (NR), styrene-butadiene rubber (SBR), and ethylene propylene diene monomer rubber (EPDM) have low microwaves absorption due to their non polar characteristic. This limitation can be overcome by the addition of conductive filler like carbon black[8-10].

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According to Scuracchio et al.[3], the technique of devulcanization by microwaves is able to generate a material with properties quite different from the original vulcanized rubber. Among the properties, the most remarkable is the ability to flow and to be remolded. This feature, allied to the possibility of its revulcanization, indicates the wide applicability of the technique. Bani et al.[11] demonstrated that microwaves can be applied easily and have many advantages, such as high heating rate, without any need of additional mechanical or chemical treatments.

On the other hand, thermoplastic vulcanized (TPVs) are a kind of polymeric blend produced via dynamic vulcanization of a dispersed elastomeric phase, i.e. the selective cross‑linking of the rubber phase while mixed with the molten thermoplastic[12]. The final morphology consists of cross-linked rubber particles dispersed in a thermoplastic matrix. The thermoplastic matrix is responsible for the processability of TPVs, while the cross-linked elastomer particles are responsible for the elasticity at room temperature.

The final morphology of this kind of material is the main responsible for the rheological and physical properties beyond being controlled by the processing conditions and characteristics of the constituting materials[13].

The balance between break-up and coalescence of the droplets of the elastomeric phase provides the final morphology of a TPV during processing. In the special case of TPV composed by devulcanized elastomer as dispersed phase, the devulcanization acts increasing the break-up ability while the revulcanization acts decreasing the coalescence, and both effects contribute to the refinement of the morphology.

In addition, higher amounts of recycled rubber can be added in the blend without properties degradation, since

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Blends of ground tire rubber devulcanized by microwaves/HDPE - Part A: influence of devulcanization process the devulcanization process tends to increase the adhesion of the particles on the thermoplastic phase.

2.4 Characterization The revulcanization characteristics of the devulcanized GTRs under different exposure times to microwaves were studied by using an oscillatory dual cone Monsanto Rheometer 100, according to ASTM D1646-07. Curves of torque versus time were obtained at 160°C. The nomenclature is type GTRX+ad, where X represents the exposure time of GTR to microwaves and “+ad” the presence of vulcanization additives.

In this work, the influence of microwaves devulcanization of the elastomeric phase in the blends GTR/HDPE is investigated. The results show that microwaves treatment of the GTR in the blends can influence the mechanical, viscoelastic, thermal and rheological properties.

2. Experimental

Thermal properties of the HDPE phase were analyzed

by Differential Scanning Calorimetry (DSC) in a DP Union DSC Q200 under nitrogen atmosphere. The samples were heated from room temperature to 190°C and were held at this temperature for 3 min to eliminate their thermal history and destroy the HDPE crystalline nuclei. They were then cooled to –90°C and were subsequently heated to 200°C. All the steps were performed at a rate of 10°C/min.

2.1 Materials HDPE IA-59, a grade for injection molding, was kindly supplied by Braskem (MFI = 7.3 g/10min). Ground waste truck tire (GTR) separated from non elastomeric components; rubber accelerator N-tert-butyl-2-benzothiazole sulfenamide (TBBS) and sulfur were kindly supplied by Pirelli Pneus Ltda.

Mechanical properties of the blends were analyzed by tensile tests in an Instron Universal Testing Machine 3369 with a 10 kN load cell at a crosshead speed of 50 mm/min. The samples were prepared in the shape of plates by compression molding at 160°C in a hydraulic press, and then the blends were cut into dumb-bell shaped tensile test according to ASTM D412, type IV.

2.2 Devulcanization of GTR and mixture with vulcanization additives GTR was devulcanized in a system comprised of a conventional microwave oven adapted with a motorized stirring system with speed control. The devulcanization process was done by using the maximum power of the oven, i.e. 820W. The time at which the material was exposed to microwaves ranged from 1 to 5 minutes and also 2-2, 2-2‑2, and 3-3, where the numbers represent the exposure time to microwaves (minutes) and the hyphen corresponds to an interval of 10 minutes between consecutive treatments, under stirring with the oven switched off.

Rheological properties of the blends were analyzed by small amplitude oscillatory rheometry in frequency sweep mode, by using a parallel plate rheometer Anton Paar CTD450 (diameter 25 mm, gap 1.3 mm, 0.5% strain for the viscoelastic linear response at 170°C under inert atmosphere).

The devulcanized GTR was mixed with the vulcanization additives by using a laboratory two roll mill PRENMAR for approximately 6 minutes at room temperature. To promote the dynamic revulcanization during the processing with the thermoplastic, 1 phr of accelerator TBBS and 1 phr of sulfur were added.

Dynamic mechanical properties of the blends were analyzed by using a Dynamic Mechanical Analyzer (DMA) Q800 TA Instruments. The analyses were performed by Single Cantilever mode, frequency of 1 Hz, temperature ranging from –100 to 140°C and heating rate of 3°C/min.

A Jeol JMS-6701F Scanning Electron Microscope was used to observe the morphology of the blends with working distance of 5.5 mm. The samples were firstly pressed by using a hydraulic press, cut, fractured just after being immersed in liquid nitrogen and then coated with golden by using a sputter coater.

2.3 Preparation of the blends The blends were prepared in an internal mixer coupled to a torque Rheometer Polylab 900 at 160°C and 80 rpm for 15 minutes. The compositions and nomenclature used for the blends are summarized in the Table 1.

Table 1. Nomenclatures and compositions of the blends produced in this work.

Nomenclature

GTR amount (wt%)

HDPE amount (wt%)

80GTR0/20HDPE 80GTR0+ad/20HDPE 80GTR1+ad/20HDPE 80GTR2+ad/20HDPE 80GTR3+ad/20HDPE 80GTR4+ad/20HDPE 80GTR5+ad/20HDPE 80GTR2-2+ad/20HDPE 80GTR2-2-2+ad/20HDPE 80GTR3-3+ad/20HDPE

80 80 80 80 80 80 80 80 80 80

20 20 20 20 20 20 20 20 20 20

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Devulcanization time of GTR (min) — — 1 2 3 4 5 2-2 2-2-2 3-3

Presence of vulcanization additives — Yes Yes Yes Yes Yes Yes Yes Yes Yes

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de Sousa, F. D. B., Gouveia, J. R., de Camargo, P. M. F., Fo., Vidotti, S. E., Scuracchio, C. H., Amurin, L. G., & Valera, T. S.

3.Results and Discussion 3.1 Revulcanization characteristics Torque versus time curves of devulcanized rubber containing vulcanization additives are showed in the Figure 1.

According to Figure 1, the samples GTR5, GTR2-2‑2 and GTR3-3 presented reversion trend, i.e. the torque measured by the equipment tends to decline at the end of the analysis. This behavior happened probably as consequence of degradation of rubber main chains, since these samples were exposed to microwaves for long periods of time.

The revulcanization characteristics of the samples, calculated from the curves presented in the Figure 1 are summarized in the Table 2. In general, it can be observed that the optimum cure time and scorch time, represented by t90 and ts1 respectively, were lower for samples with the highest exposure times to microwaves in one step as well as in multistep treatments. This behavior is characteristic of reclaimed rubber and it was observed by some other authors[14-17], which probably happens due to the presence of residual curatives from the first vulcanization. The devulcanization process increases the freedom degree of the polymeric chains, accelerating the reaction with the increasing of the exposure time of the GTR to microwaves, possibly due to the great amount of effective shocks during the process.

GTR5+ad, GTR2-2+ad and GTR3-3+ad presented lower values of ML and MH, minimum and maximum torque respectively, which demonstrates lower cross-linking

densities in comparison to the other samples. The values of the subtraction of MH-ML did not present a trend, but they were lower in relation to GTR0+ad, with exception of the samples GTR1+ad and GTR3+ad. This value is related to the cross-linking density of the sample and its reduction is attributed to the breaking of reticulation as a result of the devulcanization by microwaves.

It can be also observed that the samples with higher values of CRA (Cure Rate Average) were exposed to microwaves for 2 minutes, in two or three steps of treatment. CRA values were calculated according to Equation 1[18]:

CRA =

1 t90 − ts1

(1)

where t90 is the optimum cure time and ts1 the scorch time.

The value is proportional to the average slope of torque versus time curve or, in other words, it is proportional to the rubber revulcanization speed.

The devulcanization process reduced the ML values with the increase of the exposure time to microwaves. However, a clear trend was not observed. As the ML value is proportional to the initial viscosity of the sample[19], the increase of the exposure time to microwaves reduced the viscosity of rubber induced by the breaking of the three‑dimensional network of the vulcanized GTR. The sample GTR3 presented the highest ML value probably due to the formation of new bonds in the rubber, since during sample exposure to

Figure 1. Torque versus time curves of devulcanized rubber containing vulcanization additives. The curves were separated for better visualization and analysis of results, and GTR0 without vulcanization additives was also analyzed for comparison.

Table 2. Revulcanization behavior of the GTRs devulcanized by microwaves.

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Sample

t90 (min)

ts1 (min)

ML (dN.m)

MH (dN.m)

MH-ML (dN.m)

CRA (min–1)

GTR0+ad GTR1+ad GTR2+ad GTR3+ad GTR4+ad GTR5+ad GTR2-2+ad GTR2-2-2+ad GTR3-3+ad

2.90 2.45 1.97 2.28 2.04 1.95 1.88 2.06 1.85

2.16 1.59 1.89 1.48 1.52 1.25 1.61 1.59 1.34

27.40 24.10 18.50 29.20 22.50 16.80 15.30 22.20 14.40

34.00 30.90 20.80 38.00 28.20 23.00 17.90 27.00 17.70

6.60 6.80 2.30 8.80 5.70 6.20 2.60 4.80 3.30

1.35 1.16 12.50 1.25 1.92 1.43 3.70 2.13 1.96

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Blends of ground tire rubber devulcanized by microwaves/HDPE - Part A: influence of devulcanization process microwaves, bonds can be broken and created at the same time[20]. Therefore, depending on the time at which the sample was exposed to microwaves, the former event can exceed the later one or vice-versa. On the other hand, the sample GTR3-3 presented the lowest ML value obtained, which may be due to degradation of the rubber by the high exposure time of the sample to microwaves.

MFinal. The fluidity of rubber facilitates the processing, its dispersion into other polymers to form a polymer blend, as well as the revulcanization reaction.

3.3 Oscillatory rheometry The storage modulus (G’) and complex viscosity (η*) of the blends 80GTR/20HDPE, as function of the frequency, are summarized in the Figure 2.

3.2 Processing behavior of the blends

In order to facilitate the analysis of the results, a table of the G’ at the minimum and maximum frequencies of all the blends was created (Table 4).

According to Shahbikian et al.[21], the advantage of using internal mixer to produce TPV is the possibility of monitoring the effect of each component on the torque/temperature evolution of the blend. Some researchers use this advantage to examine different phenomena such as dynamic vulcanization[21-29], as performed in this work.

During the mixing, just after the addition of the matrix phase and as soon as the torque measured by the equipment was stabilized, GTR (containing or not vulcanization additives) was added into the mixer, what permitted the analysis of the dynamic revulcanization behavior of the blends, which is shown in Table 3. The MFinal values represent the torque measured by the equipment at the end of the mixing process.

In general, t90 and ts1 values of the blends were much smaller than the values of the neat rubber obtained by using a rheometer (Table 2), which shows that the dynamic revulcanization reaction occurred with higher rate. Consequently, the CRA values of the blends were also higher in comparison to the neat rubber, which confirms, as just verified through the results of t90 and ts1, that the dynamic revulcanization reaction occurred more quickly in comparison to revulcanization (exception: GTR2+ad, in which the highest reaction rate happened). It happened possibly due to higher shear rates generated within the internal mixer during processing.

However, the values of ML, MH and the subtraction of MH-ML were not analyzed in this section and compared with the ones of the neat rubber (section 3.1), since they also take into consideration the viscosity of the HDPE phase (among other factors) and it may lead to erroneous conclusions.

In general, but with some exceptions, it was verified a trend towards the reduction in the MFinal values of blends (concerning the final viscosity of blends) as the exposure time of GTR to microwaves got higher, what demonstrates that the GTR devulcanization increased the fluidity of this phase. The cross-linking density and a possible degradation of the thermoplastic phase may also have influenced the

According to the Figure 2, the complex viscosity decreased with the increase of the frequency, which clearly shows the pseudoplastic behavior of the blends, assuming the Cox Merz rule[30-35]. η* of the dynamically revulcanized blends are higher than the blend 80GTR0/20HDPE due the increase of the cross-linking density of the GTR phase[35].

Being G’ proportional to the stored energy[36], this value is proportional to the elasticity or, in other words, to the cross-linking density of the elastomeric phase of the blend.

G’ is influenced also by morphology of the blends[28,37].

The morphology refinement and compatibility tend to increase the G’ values. According to SEM micrographs, no conclusion about the morphology refinement of the blends can be made, but the mechanical properties results alert to the poor adhesion between the phases. The blends that presented the lowest elongation at break results were the same that presented the lowest G’ at minimum frequency (80GTR4+ad/20HDPE, 80GTR5+ad/20HDPE and 80GTR3‑3+ad/20HDPE), which can have been result of the poor adhesion, occurring a possible particle detachment from the matrix when applied an external stress.

3.4 Dynamic mechanical properties The temperature dependence of tan δ of the blends is shown in the Figure 3.

According to the Figure 3, there are two transitions related to the phases of the blends: the first one around -30°C refers to the glass transition (Tg) of the GTR and the other refers to α transition of the HDPE phase (Tα) around 100°C. The existence of two distinct transitions confirms the immiscible character of the blends. It can also be observed that there is a trend towards the reduction of the area under the peak related to GTR transition, as well as the reduction

Table 3. Dynamic revulcanization behavior of the of the blends 80GTR/20HDPE. Blend

t90 (min)

ts1 (min)

CRA (min–1)

MFinal (dN.m)

80GTR0/20HDPE 80GTR0+ad/20HDPE 80GTR1+ad/20HDPE 80GTR2+ad/20HDPE 80GTR3+ad/20HDPE 80GTR4+ad/20HDPE 80GTR5+ad/20HDPE 80GTR2-2+ad/20HDPE 80GTR2-2-2+ad/20HDPE 80GTR3-3+ad/20HDPE

1.17 0.95 0.65 1.15 1.13 1.01 0.80 0.75 1.05

0.75 0.75 0.48 0.95 0.72 0.68 0.45 0.50 0.68

2.38 5.00 5.88 5.00 2.44 3.03 2.86 4.00 2.70

107.00 108.00 106.00 111.00 92.50 93.90 63.10 103.00 96.20 54.30

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de Sousa, F. D. B., Gouveia, J. R., de Camargo, P. M. F., Fo., Vidotti, S. E., Scuracchio, C. H., Amurin, L. G., & Valera, T. S. of the height of the same peak, which is due to mobility restriction generated by cross-linkings of this phase[32,38,39].

Tg values of the devulcanized rubber and rubber phases of the blends were obtained from the values of maximum peaks of the curves tan δ versus temperature. These values are shown in the Figure 4.

According to the Figure 4, three zones of distinct Tg behaviors can be determined. They were divided into continuous, dotted and dashed line zones, which are described below.

Table 4. G’ at the minimum (0.01 rad/s) and maximum (300 rad/s) frequencies of the blends.

Blend 80GTR0/20HDPE 80GTR0+ad/20HDPE 80GTR1+ad/20HDPE 80GTR2+ad/20HDPE 80GTR3+ad/20HDPE 80GTR4+ad/20HDPE 80GTR5+ad/20HDPE 80GTR2-2+ad/20HDPE 80GTR2-2-2+ad/20HDPE 80GTR3-3+ad/20HDPE

G’ (Pa) at 0.01 rad/s 1.52x105 2.16x105 1.82x105 1.76x105 2.34x105 1.98x105 1.75x105 2.27x105 2.34x105 2.22x105

G’ (Pa) at 300 rad/s 5.75x105 5.52x105 5.19x105 4.95x105 6.69x105 5.82x105 5.84x105 5.87x105 6.69x105 6.73x105

Continuous line zone: the GTR was not exposed to microwaves.

Dotted line zones: the final temperature of the GTRs after the exposure time to microwaves probably was not enough to provide high degree of devulcanization in the samples.

Due to the low degree of devulcanization, there was not a significant change in Tg of the rubber, which behaved just like a vulcanized one.

Dashed line zones: the final temperature of the GTRs after the exposure time to microwaves was enough to generate high degree of devulcanization in the sample. During processing of the blends, due to the devulcanization degree reached by the elastomeric phase of the samples, the rubber chains acquired some mobility, demonstrated by the increase in the Tg values. In other words, the devulcanization level of the elastomeric phase influenced the dynamic revulcanization reaction, changing the Tg value of this phase.

3.5 Thermal properties by DSC The results of the DSC obtained from the second heating cycle of the blends are shown in the Table 5. The crystallization degree was calculated according to Equation 2[40]:  ∆H  χc = m .100 (2) ( ∆H m100 .WHDPE )  

Figure 2. G’ and η* versus frequency of the blends. The curves were separated for better visualization and analysis of results.

Figure 3. Tan δ versus temperature of the blends 80GTR/20HDPE. The curves were separated for better visualization and analysis of the results.

260

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Blends of ground tire rubber devulcanized by microwaves/HDPE - Part A: influence of devulcanization process where χc is the crystallization degree, ΔHm is the enthalpy of melting (J/g), ΔHm100 is the enthalpy of melting of the HDPE 100% crystalline (293 J/g)[41] and WHDPE is the mass fraction of HDPE in blend.

The melting temperatures of the HDPE phase did not present large variations in function of the exposure time of the GRT to microwaves. However, the crystallization degree of the HDPE phase was affected by the presence of the rubber phase.

The blends that presented the highest crystallization degree of the HDPE phase were the ones in which the GTR phase was exposed to microwaves for longer periods of time (with exception of the samples 80GTR0+ad/20HDPE and 80GTR3+ad/20HDPE). This fact is probably due to

the more pronounced refinement of the morphology of the blends with the increase of the exposure of the GTR to microwaves, which could not be observed by SEM micrographs. According to Utracki (p. 248)[42], “[...] the finer the amorphous droplets are dispersed, the larger the total interfacial contact surface, and thus the higher is the possibility of nucleation at these interfaces.”

3.6 SEM The morphologies of some blends are shown in the Figure 5.

According to the SEM micrographs, it could be observed that the blends containing GTR with longer exposure time to microwaves presented a less coarse surface in comparison to the other blends as a result of a lower fracture resistance to the external force applied on the blends. This tendency was also observed in the results of mechanical properties (see section 3.7). Regarding the morphology refinement, no conclusions can be made because the impossibility to distinguish the phases from the presented SEM micrographs.

In the blend 80GTR3-3/20HDPE some voids can clearly be observed (arrows in the Figure 5f) probably due to the degradation of the rubber phase and poor degree of interfacial adhesion between the phases of the blend.

3.7 Mechanical properties

Figure 4. Tg values of the devulcanized rubber and rubber phases of the blends as determined by DMA.

The main results of the tensile tests of the blends are presented in the Table 6.

On the whole but with some exceptions, the values of stress at break and elongation at break reduced as the exposure time of GTR to microwaves got higher, while the values of Young’s modulus presented an opposite behavior. The tensile

Figure 5. SEM micrographs of the blends: (a) 80GTR0/20HDPE, (b) 80GTR0+ad/20HDPE, (c) 80GTR3+ad/20HDPE, (d) 80GTR4+ad/20HDPE, (e) 80GTR2-2+ad/20HDPE, (f) 80GTR3-3+ad/20HDPE. Polímeros, 25(3), 256-264, 2015

261

de Sousa, F. D. B., Gouveia, J. R., de Camargo, P. M. F., Fo., Vidotti, S. E., Scuracchio, C. H., Amurin, L. G., & Valera, T. S. Table 5. Values of melting temperature, enthalpy of melting (ΔHm) and crystallization degree (χc) of the HDPE phase of the blends.

Sample HDPE 80GTR0/20HDPE 80GTR0+ad/20HDPE 80GTR1+ad/20HDPE 80GTR2+ad/20HDPE 80GTR3+ad/20HDPE

Tm (oC) 141.74 134.86 132.12 133.93 133.53 133.14

ΔHm (J/g) 183.72 52.84 33.57 41.38 40.43 51.29

χC (%) 62.70 90.17 57.29 70.61 68.99 87.52

Sample

Tm (oC)

ΔHm (J/g)

χC (%)

80GTR4+ad/20HDPE 80GTR5+ad/20HDPE 80GTR2-2+ad/20HDPE 80GTR2-2-2+ad/20HDPE 80GTR3-3+ad/20HDPE

133.81 135.79 133.46 136.65 133.77

39.02 52.03 51.75 54.56 46.66

66.59 88.79 88.30 93.10 79.63

Table 6. Mechanical properties of the blends.

Blend 80GTR0/20HDPE 80GTR0+ad/20HDPE 80GTR1+ad/20HDPE 80GTR2+ad/20HDPE 80GTR3+ad/20HDPE 80GTR4+ad/20HDPE 80GTR5+ad/20HDPE 80GTR2-2+ad/20HDPE 80GTR2-2-2+ad/20HDPE 80GTR3-3+ad/20HDPE

Young’s modulus (MPa) 36.02±7.97 33.33±1.99 27.98±0.67 30.60±1.98 33.94±3.95 37.30±3.54 39.49±1.33 25.01±0.86 48.70±2.81 57.02±3.37

strength of the blends did not vary significantly with the increase of the exposure time of the GTR to microwaves.

According to Prut et al.[36], Young’s modulus depends on the crystallinity development during the quenching.

With the increase of the crystallinity, the matrix became tougher, which led to the increase of the Young’s modulus.

The increase of the crystallinity of the matrix could also generate less perfect crystals, which also may have resulted into an increase of the Young’s modulus values.

Another observation is that, in general, the results obtained were not so good, especially the results of elongation at break, for blends which are supposed to be TPVs. According to Hong and Isayev[19], adhesion between the GRT and polymer matrix is one of the major factors controlling the mechanical properties of such blends. Also, according to some authors[43,44], deterioration on the elongation at break is due to the poor interfacial adhesion between the phases.

According to the presented results, adhesion between the phases was not sufficient to promote good stress transference.

The blends containing GTR0 and vulcanization additives presented higher mechanical properties in comparison to the same ones without additives. These results showed that the dynamic revulcanization improves the mechanical properties, as also observed by other authors[39,45-51].

Luo and Isayev[46] studied the properties of the blends polypropylene (PP)/GTR devulcanized by ultrasound using different curing systems and processing routes, and all the blends presented low elongation at break. Some other authors[19,52] also achieved the same results. The reason for this behavior is the large size of rubber particles and the premature curing of this phase when the curatives are poorly distributed in the rubber[46]. Also, according to Antunes et al.

[22] , when the dynamic vulcanization happens, the curatives are not well distributed by the batch mixer, resulting in different 262

Stress at break (MPa) 60.61±5.45 92.93±7.30 54.87±8.36 80.59±6.56 79.98±5.02 54.48±8.35 64.36±15.95 79.25±2.91 53.63±10.62 40.67±4.70

Tensile strength (MPa) 2.72±0.24 4.30±0.34 3.01±0.46 3.67±0.34 3.73±0.19 2.99±0.46 3.53±0.87 4.26±0.94 2.94±0.58 2.23±0.18

Elongation at break (mm/mm) 0.40±0.06 0.59±0.10 0.25±0.06 0.53±0.10 0.40±0.06 0.15±0.05 0.16±0.05 0.47±0.06 0.19±0.05 0.09±0.01

levels of cross-linkings. In the present work, this problem was avoided, since the vulcanization agents were previously added in the rubber phase and mixed by using a two roll mill. So, the reason for the poor mechanical properties is probably the lack of adhesion between the phases.

One of the qualifying standards for a blend to be deemed as a TPV is to present typical elastomeric elongation, which has also not been verified in the obtained results. However, these blends have high concentrations of GTR phase (80% in mass), a recycled material, what may have deteriorated the mechanical properties. Grigoryeva et al.[53] produced dynamically vulcanized blends using GTR, and in some of the production methods used by the authors, TPVs were also not obtained. According to the authors, in these cases there was not an effective interfacial stress transference between the phases, and not an entanglement of the GTR rubber chains into the surrounding matrix. These facts could also have happened in this work. Therefore, additional studies must be performed, taking into consideration the use of compatibilizer additives or a way to improve the interfacial characteristics of the blends like using nanofillers.

4. Conclusions The dynamically revulcanized blends based on GTR devulcanized by microwaves (under different exposure times) and HDPE were analyzed by different techniques.

According to the torque development during the mixing process, dynamic revulcanization was faster than the revulcanization of the neat rubber, due to the high shear rates generated during the processing. The oscillatory rheometry results showed that the lack of adhesion between the phases influenced the rheological properties of the blends, which resulted into poor mechanical properties, especially in the blends containing GTR exposed to microwaves for longer Polímeros , 25(3), 256-264, 2015

Blends of ground tire rubber devulcanized by microwaves/HDPE - Part A: influence of devulcanization process exposure times. The dynamic mechanical properties showed that there were differences in the Tg values of the elastomeric phase, depending on the exposure time to microwaves.

However, no conclusion about the morphology refinement of the blends can be made based on the SEM micrographs.

Summarizing, devulcanization process of GTR can change completely the final properties of the revulcanized blends 80GTR/20HDPE, since it changes the fluidity of rubber during processing. The process parameters like exposure time can be analyzed based on the final properties of these blends. Devulcanization by microwaves can be a strong alternative to solve the problem of disposal of waste rubber.

5. Acknowledgements The authors would like to thank Braskem and Pirelli for the material donation, Departamento de Engenharia de Materiais - DEMa, Universidade Federal de São Carlos for the laboratory facilities, FAPESP (process number 2010/15799-6) and CNPq (process number 201891/2011-5) for the financial support.

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Received: Apr. 14, 2014 Revised: Oct. 07, 2014 Accepted: Nov. 11, 2014 Polímeros , 25(3), 256-264, 2015

http://dx.doi.org/10.1590/0104-1428.1818

Immobilization of myoglobin in sodium alginate composite membranes Katia Cecília de Souza Figueiredo1, Wilbert van de Ven2, Matthias Wessling2, Tito Lívio Moitinho Alves1 and Cristiano Piacsek Borges1 1 Instituto Alberto Luiz Coimbra de Pós-Graduação e Pesquisa de Engenharia – COPPE, Universidade Federal do Rio de Janeiro – UFRJ, CEP 21941-914, Rio de Janeiro, RJ, Brazil 2 Membrane Technology Group, Science and Technology, University of Twente – UT, Meander P. O. Box 217, 7500, AE Enschede, The Netherlands

*katia@deq.ufmg.br

Abstract The immobilization of myoglobin in sodium alginate films was investigated with the aim of evaluating the protein stability in an ionic polymeric matrix. Myoglobin was chosen due to the resemblance to each hemoglobin tetramer.

Sodium alginate, being a natural polysaccharide, was selected as the polymeric matrix because of its chemical structure and film-forming ability. To improve the mechanical resistance of sodium alginate films, the polymer was deposited over the surface of a cellulose acetate support by means of ultrafiltration. The ionic crosslink of sodium alginate was investigated by calcium ions. Composite membrane characterization comprised water swelling tests, water flux, SEM images and UV-visible spectroscopy. The electrostatic interaction between the protein and the polysaccharide did not damage the UV-visible pattern of native myoglobin. A good affinity between sodium alginate and cellulose acetate was observed. The top layer of the dense composite membrane successfully immobilized Myoglobin, retaining the native UV-visible pattern for two months.

Keywords: myoglobin, sodium alginate, immobilization, composite membrane, cellulose acetate.

1. Introduction Facilitated transport membranes have been investigated as an alternative to increase flux and selectivity of ordinary films[1]. These special properties are achieved by adding a carrier to the membrane, so that the desired component is preferentially transported from the feed to the permeate stream. The reaction between the solute and the carrier must be reversible. The process can be described as follows: 1) the solute binds to the carrier on the feed/membrane interface, 2) the complex diffuses through the membrane due to a concentration gradient, 3) the solute is released on the membrane/permeate interface and the carrier is regenerated.

The transport mechanism for fixed site carriers is similar to the mobile ones, but the solute should be able to jump from one site to the other until it reaches the membrane/permeate interface, where it is released. Due to this reactive mechanism, the effective transport of solutes can only be achieved for low partial pressure of the desired component, otherwise it would cause the saturation of the carrier sites. Therefore, one of the most promising applications of such membranes is in the air separation, due to the similar physical properties of its components. The high demand for pure oxygen and oxygen-enriched air has motivated a large number of studies on oxygen facilitated transport membranes[2].

The use of the synthetic oxygen carriers in supported liquid membranes was extensively studied in the last 30 years and the results were remarkable to the understanding of the solute transport mechanism through the membrane.

However, the lack of stability in long-term applications due to the evaporative loss of solvent and oxidation of the

Polímeros, 25(3), 265-270, 2015

carrier limited the advance of this approach. Alternatively, new morphologies have been tested, as the encapsulation of the carrier in a polymeric shell[3] or the use of a polymeric matrix to stabilize the carrier[4-6]. High selectivities have been reported but the instability of such membranes remains unsolved and there is still a need for the development of an efficient process to replace the energy-intense cryogenic distillation.

The regained interest in the use of the hemoproteins in oxygen facilitated transport membranes relies on the role of the globin moiety to prevent the autoxidation of these biocarriers[7] and the development of the molecular biology tools, which allows the change in the aminoacids sequence to provide more stable oxygen carriers[8]. The choice of myoglobin to produce oxygen facilitated transport membranes was based on the fact that the protein is a single-site oxygen biocarrier and, thus, simpler than hemoglobin structure.

Sodium alginate was chosen as the polymeric matrix for the immobilization of myoglobin because of its industrial and upcoming biotechnological applications as thickening and emulsifying agent used in food formulations. Also it suffers gelation in a quite gentle way in the presence of ions such as calcium.

Being an anionic copolymer of 1,4-linked-β-D-mannuronic acid (M) and α-L-guluronic acid (G) residues, alginate forms a gel due to the stacking of guluronic acid blocks, which have a particular diamond shape, ideal for the cooperative binding of calcium and other ions[9]. The selective ion binding characteristic, which forms a structure well-known

265

S S S S S S S S S S S S S S S S S S S S

Figueiredo, K. C. S., van de Ven, W., Wessling, M., Alves, T. L. M., & Borges, C. P as the egg-box model, makes possible the immobilization of biomolecules and living cells in this polysaccharide, by tailoring the water-swelling behavior of the gel. This potential feature was exploited here aiming the stabilization of the protein inside the solid membrane. In addition to it, it is also worthwhile to consider that recent research in sodium alginate ultrafiltration shows that this polyelectrolyte has different conformations in aqueous solution as a function of the ionic environment[10]. Therefore, regarding to the synthesis of a dense composite membrane, the top layer density could be controlled by adding non-gelling ions to the polymeric aqueous solution.

In this work, the immobilization of myoglobin in a sodium alginate (SA) matrix was investigated aiming the stabilization of the protein in the polymeric membrane.

In order to circumvent the lack of mechanical strength of SA films, composite dense membranes were tested. The use of a thin active polymeric layer in which both the oxygen uptake and release can be performed in one single stage will improve process efficiency and economy[11].

cellulose acetate substrate (MWCO 100 kDa) in order to prepare composite membranes. The choice of the ultrafiltration cellulose acetate support was based on its similarity with sodium alginate. The porous support was prior characterized by the ultrafiltration of pure water, alginate (0.1 wt%) and myoglobin (0.1 wt %) aqueous solutions, at 20oC, feed to permeate pressure difference of 3 bar and 700 rpm.

Composite membrane preparation was based on the pore blockage of the cellulose acetate support by means of the ultrafiltration of SA aqueous solution (0.1 wt%) containing 10 mM KCl. The addition of the saline solution can decrease the space between the anionic copolymer. The experimental set up was composed by a dead-end filtration tank with 36.3 cm3 of effective area, in which the SA/KCl aqueous solution was added. The UF support was placed on the bottom of the tank. Once the pores were blocked, which was measured by the reduction on permeate flux, oxymyoglobin or the crosslinking ion (Ca2+ or Fe2+ aqueous solution) was added to the filtration tank. Composite membranes were then dried at 20oC overnight and stored in a desiccator.

2. Experimental

2.5. Membrane characterization

2.1. Materials

2.5.1. Integral membranes

Myoglobin from horse heart (Fluka), minimum 90% of purity, was obtained in the metmyoglobin form (Fe3+). Sodium alginate medium viscosity and Sephadex G25 were acquired from Aldrich. KCl, CaCl2, FeCl2 and Na2S2O4 were used as received. Centripep device (Millipore, Bedford, MA) and ultrafiltration cellulose acetate substrate (MWCO 100 kDa, Nadir) were used in protein concentration and as a porous support on the composite membrane.

Membrane thickness was determined by means of a micrometer. The result was the average of at least 20 measures.

Myoglobin physiological form in the dense film was investigated in an UV-visible spectroscope (UV mini 1240, Shimadzu, Tokyo, Japan). Strips of (0.4 x 2) cm2 of dense membranes were placed in the cuvette. SA films with no added-myoglobin were used as the reference. This technique was used qualitative in order to evaluate the iron oxidation state of the immobilized myoglobin film, since Fe2+ and Fe3+ show different UV-visible absorption spectra. For instance, the Soret peak is 409 nm for Fe3+, while it is at 418 nm for Fe2+ bound to oxygen.

Water swelling degree tests were conducted to characterize the membranes. Membrane strips were placed in a Petri dish with water, at 20oC. They were removed 48 hours later, had their weight determined and were dried in a vacuum oven, at 30oC. The swelling degree (SD) was calculated according to Equation 1, in which mwet and mdry are the weight of wet and dried membrane, respectively.

2.2. Myoglobin activation Metmyoglobin (Fe3+) was activated, i. e., reduced to Fe2+, by reacting the protein aqueous solution and sodium dithionite. Oxymyoglobin (Fe2+ bound to O2) was separated from metmyoglobin by means of a chromatographic method using Sephadex G25-filled column at room temperature.

Once the fraction was purified, its concentration was adjusted by means of centrifugation on a Centripep device using an ultrafiltration filter with MWCO of 5 kDa, at 1,800 g and 4oC.

2.3. Integral membrane preparation Sodium alginate, SA, and KCl were dissolved in distilled water in room temperature. Myoglobin was added to the solution in the desired amount. The mixture was gently stirred in order to dissolve the macromolecule with no damage to the structure. The mass ratio between Mb and SA was varied from 0 to 2 g/g. The mixture was poured into a Petri dish and dried at 20oC, for 15 hours. This method is known as solvent evaporation. The final contents of SA and KCl on the casting solution were 1.0 wt% and 10 mM, respectively.

The pH of the casting solution was 6.5. Membranes were pelled of the plate and stored in a desiccator.

2.4. Composite membrane preparation The preparation of composite membranes was performed with the aim of increasing the mechanical strength of SA films. Sodium alginate was deposited over an ultrafiltration 266

SD =

( mwet − mdry ) *100 (1) mdry

2.5.2. Composite membranes The characterization of composite SA membranes comprised vapor water swelling tests, pure water flux and scanning electron microscopy.

Vapor water swelling degree of the integral and composite SA membrane was evaluated by placing the samples in a closed recipient containing liquid water in the bottom, as described elsewhere[12]. Samples were cut in (2 x 2) cm2, had their dry weight measured and were placed in a petri dish, which was transferred to the recipient, which was closed for 48 hours, at 20°C. After that, they were removed from the recipient and had their wet weight determined. The swelling degree (SD) was calculated according to Equation 1.

Polímeros , 25(3), 265-270, 2015

Immobilization of myoglobin in sodium alginate composite membranes Pure water flux through the composite SA membranes was determined in filtration experiments conducted in a 3.14 cm2 area cell. Milli-Q water was filtrated at 3 bar, 20 °C and 700 rpm. The permeate weight was measured in time intervals of 15 minutes and the corresponding flux was calculated.

The composite membrane morphology was characterized by scanning electron microscopy, SEM (JEOL, Tokyo, Japan).

Samples were freeze-dried in liquid nitrogen, cut and coated by gold sputtering of the surface.

In addition to it, the amounts of alginate and myoglobin retained by the UF support were investigated by spectroscopic analysis. On the first case, the polysaccharide was quantified by the Dubois technique[13], while the protein quantification was conducted by means of the absorption on the Soret region, by the coefficient of extinction[14].

3. Results and Discussion 3.1. Integral membranes The addition of KCl was investigated in order to assure a denser layer, since the polysaccharide chains are less extended in the presence of such non-gelling ions[10].

Red and defect‑free membranes were obtained. The average thickness of the films was 30 μm. Once the pH of the casting solution was 6.5, SA was negatively charged while myoglobin was positively charged, since the protein pI ranges from 7.0 to 7.4[15]. These data suggested that the main interaction between the polysaccharide and the protein was electrostatic.

As the electrostatic is a strong interaction, it could lead to protein denaturation, which is usually also associated with decrease in the ability to carry oxygen. On the other hand, the electrostatic interaction may act as a stabilizer of the ionic charged aminoacids that composes the macromolecule.

This dubious behavior was investigated by means of an UV-visible spectroscopy of the films. Being a hemoprotein macromolecule, myoglobin presents a specific UV-visible pattern, which is frequently used in order to infer the physiological form of this protein. Such data is able to inform about the iron microenvironment, but it is related to the prosthetic group of the protein[16].

The spectroscopic analysis of the dense membrane is presented in Figure 1. Membranes with no protein were used as reference. The absorption peaks at 418 (Soret), 542 and 581 nm were equivalent to the oxymyoglobin UV-visible pattern[15] and showed that the protein active site (Soret region) was retained in the polysaccharide matrix. The absorbance scan of these membranes was repeated once a week for two months and no change on the standard pattern was noticed.

These results indicated that the immobilization of the protein in a saline polysaccharide film did not show changes in the primary structure of the macromolecule. A detailed study should be carried in order to evaluate the protein tertiary and quaternary structures, by means of circular dichroism, for instance.

Water swelling degree tests of integral dense membranes were conducted to evaluate the interaction of SA and Mb. As myoglobin is positively charged in the pH of 6.5, it could work as an ionic crosslinker for SA membranes.

However, the membranes had high swelling degree, with Polímeros, 25(3), 265-270, 2015

Figure 1. UV-visible spectra of SA integral membrane containing Mb. (a) Mb/SA = 2 g/g, (b) Mb/SA = 1 g/g.

lose of its shape due to the partial dissolution of the strips.

These results showed that the interaction of the polymer and protein with water was stronger than between them.

Therefore, myoglobin was considered a poor crosslinker for alginate molecules. This can be reasoned in terms of protein shape, which is a globular one, and consequently hard to act as crosslinker.

3.2. Composite membranes Pure water flux through the commercial ultrafiltration support was (38 + 2) x 10 L/h.m2, in accordance with the range expected from the manufacturer, from 200 to 400 L/h.m2. Operating conditions were 20°C, 3 bar, 700 rpm. The characterization of SA and myoglobin fluxes through the UF support can be visualized in Figures 2a and b, respectively. The fluxes through the substrate decreased with time, denoting that the pores were blocked by the solutes, which characterizes support fouling. Regarding the sodium alginate aqueous solution (Figure 2a), the triplicates showed low deviation.

The decrease on the permeate flux was ascribed to the increase on the fouling layer by pore blockage mechanism, which increased the mass transfer resistance. A similar system was investigated in order to allow the ultrafiltration of proteins with similar molecular weight and the results showed that membrane selectivity was significantly changed by pH and ionic strength of the media[17].

The amount of protein and sodium alginate retained on the substrate were addressed by spectroscopic measurements of the feed and permeate streams. More than 95% of the polysaccharide (SA) were retained on the top of the cellulose acetate ultrafiltration support, which represented 1.67 mg SA/cm2 of substrate. Regarding the protein, it was observed that 65% of myoglobin was retained on the ultrafiltration support due to the hydrophilic interaction between this positively charged solute and the porous support. Due to the high value-added protein, this simple deposition method of Mb over the support was not used, considering the macromolecule loss of 25%. It is worth to notice that Mb UV-visible spectra in the aqueous permeate were maintained, which indicates that Mb native primary structure was maintained even at a transmembrane pressure of 3 bar.

267

Figueiredo, K. C. S., van de Ven, W., Wessling, M., Alves, T. L. M., & Borges, C. P The addition of calcium chloride to the sodium alginate composite membranes was investigated aiming to crosslink the polysaccharide. This feature could be interesting in the case of investigation of oxygen transport in aqueous medium, as in the development of an artificial gill. Therefore, the amount of calcium chloride was varied and the results for the water vapor swelling are presented in Table 1. Membranes were prepared in the absence of Mb. The SA membrane with no calcium had the highest water swelling degree, as expected. Compared to the support without coating, there was an increase of 7-fold on the swelling degree, probably due to the hydrophilic sodium alginate layer. The addition of calcium chloride decreased the swelling degree, which can be reasoned in terms of the crosslinking effect of calcium ions, which were probably placed among the guluronic blocks, decreasing the macromolecular chains spacing.

It was noticed that the swelling degree was lower with the Table 1. Vapor water swelling behavior of SA composite membranes.

Sample No coated support Support + SA (no crosslinker) Support + Ca2+/SA 4 g/g Support + Ca2+/SA 16 g/g

SD (%) 6.0 ± 0.5 42.8 ± 0.7 10.6 ± 0.5 7.6 ± 0.8

increase in the calcium amount, as expected. Tests were performed in triplicates and the experimental fluctuation was around 10%.

Water permeation tests showed that no significant amount of the solvent was permeated through the crosslinked membranes during 7 hours of experiment, at 3 bar. This could be explained due to the high crosslink density of the deposited layer, which decreased alginate hydrophilic character and water flux through the membrane. This result showed that the membrane was non-porous.

SEM pictures of the membranes are depicted in Figure 3. The structure of the support without coating layer is presented in Figure 3a. It showed the difference on the layers spacing from the top to the bottom. It was noticed a good interaction between the alginate layer and the support material, due to the delamination of the top layer from the non-woven material, as can be clearly seen in Figure 3b.

The increase in calcium content led to the brittleness of the top layer, causing a defect on membrane surface, as can be seen in Figure 3c.

The study of myoglobin stability in ionic solutions, namely, CaCl2 and FeCl2, was accomplished in order to select the suitable conditions to perform the membrane crosslinking.

The use of iron (II) as an alternative to calcium was based on its presence on myoglobin prosthetic group. The spectra for

Figure 2. SA (a) and Mb (b) aqueous solution flux through cellulose acetate UF support.

Figure 3. SEM pictures of SA composite membranes. (a) Cross section of the support without coating, (b) Delamination of SA top layer, (c) Surface of Ca2+/SA 16 g/g membrane.

268

Polímeros , 25(3), 265-270, 2015

Immobilization of myoglobin in sodium alginate composite membranes

Figure 4. (a) Effect of Ca2+ on Mb structure 1. a) CaCl2 1 μM, b) CaCl2 0.07 mM, c) CaCl2 0.13 mM, d) CaCl2 0.25 mM, e) CaCl2 0.5 mM, f) CaCl2 1 mM. (b) Effect of Fe2+ on Mb structure. a) FeCl2 0.07 mM, b) FeCl2 0.13 mM, c) FeCl2 0.25 mM, d) FeCl2 0.5 mM, e) FeCl2 1 mM.

myoglobin solutions containing calcium and iron chloride at 20°C are presented on Figures 4a and b. The oxidation of myoglobin (change in iron II to iron III state) was noticed in both cases, since the two characteristic peaks of oxymyoglobin (542 and 581 nm) were replaced by a broad peak at 630 nm, typical of metmyoglobin form. However, it was noticed that the iron (II) also damaged the protein primary structure, since the spectrum for 1 mM FeCl2 solution was quite different from the standard. It was shown an increase on the absorption at the Soret region, typical of the exposition of the prosthetic group to the medium. As a result, only diluted calcium chloride solutions were investigated as crosslinker.

The use of calcium chloride solution as crosslinking agent was restricted to the concentration of 1 μM. However, the addition of calcium ions also lead to damage on myoglobin primary structure, no matter its addition was made before or after the protein solution. The reactivation of myoglobin by sodium dithionite was no longer accomplished. Membrane color changed from brown to green, which indicated some kind of chemical reaction of iron.

As an alternative, the synthesis of membrane without crosslinker was performed, varying the amount of myoglobin added to the SA top layer. This material could be used in air fractionation, aiming the increase in oxygen flux and selectivity against nitrogen. However, the use of such membrane in aqueous media is not suitable, since it showed high swelling degree. Tests revealed that the optimum mass ratio between myoglobin and SA is 0.5 g/g. Higher protein/polysaccharide ratio lead to the clustering of protein molecules.

Gas permeation tests of the Mb/SA membranes were conducted for oxygen and nitrogen, but the permeability was lower than 0.05 Barrer, the lower value detectable by our equipment. In order to prepare a more permeable system, we believe that the use of different crosslinking agents should be investigated, as dialdehydes, which could also increase the protein load in the composite membranes.

Polímeros, 25(3), 265-270, 2015

4. Conclusion The immobilization of myoglobin in sodium alginate film was successfully accomplished in this study since protein primary structure was maintained for at least two months.

The electrostatic interaction between the protein and the polysaccharide did not damage the Soret microenvironment of the biocarrier structure. This result indicated that sodium alginate is a suitable polymer for the immobilization of myoglobin.

The preparation of the composite membrane to improve the mechanical resistance of Mb-containing SA films showed high affinity between sodium alginate and the cellulose acetate UF support. The retention of the polysaccharide was higher than 95%, with ca. 1.67 mg SA/cm2. The procedure to produce non-crosslinked myoglobin-containing sodium alginate membranes was established. These results showed homogeneous membranes in which the primary native folding of myoglobin was maintained. The mass ratio between myoglobin and polymer was set in 0.5 g/g. Higher content of Mb caused protein clustering.

5. Acknowledgements Financial support from the National Council for Scientific and Technological Development (CNPq) and the Brazilian Federal Agency for Support and Evaluation of Graduate Education (CAPES Foundation) is gratefully acknowledged. Katia C. S. Figueiredo is thankful to Membrane Technology Group, in University of Twente, The Netherlands.

6. References 1. Baker, R. W. (2004). Membrane technology and applications.

London: John Wiley.

2. Baker, R. W. (2002). Future directions of membrane gas separation technology. Industrial & Engineering Chemistry 269

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3. Figoli, A., Sager, W. F. C., & Mulder, M. H. V. (2001).

Facilitated oxygen transport in liquid membranes: review and new concepts. Journal of Membrane Science, 181(1), 97-110. http://dx.doi.org/10.1016/S0376-7388(00)00508-1.

4. Shentu, B., & Nishide, H. (2003). Facilitated oxygen transport membranes of picket-fence cobaltporphyrin complexed with various polymer matrixes. Industrial & Engineering Chemistry Research, 42(24), 5954-5958. http://dx.doi.org/10.1021/ ie020770e.

5. Suzuki, T., Yasuda, H., Nishide, H., Chen, X. S., & Tsuchida, E. (1996). Electrochemical measurement of facilitated oxygen transport through a polymer membrane containing cobaltporphyrin as a fixed carrier. Journal of Membrane Science, 112(2), 155160. http://dx.doi.org/10.1016/0376-7388(95)00291-X. 6. Nishide, H., Tsukahara, Y., & Tsuchida, E. (1998). Highly selective oxygen permeation through a poly(vinylidene dichloride)-cobalt porphyrin membrane: hopping transport of oxygen via the fixed cobalt porphyrin carrier. The Journal of Physical Chemistry B, 102(44), 8766-8770. http://dx.doi. org/10.1021/jp9816317.

7. Sugawara, Y., Matsuoka, A., Kaino, A., & Shikama, K. (1995). Role of globin moiety in the autoxidation reaction of oxymyoglobin: effect of 8 M urea. Biophysical, 69(2), 583-592. http://dx.doi.org/10.1016/S0006-3495(95)79932-5.

PMid:8527673.

8. Ferraz, H. C. (2003). Membranas de transporte facilitado para separação de oxigênio utilizando biotransportadores (Doctor Thesis). Universidade Federal do Rio de Janeiro, Rio de Janeiro.

9. Draget, K. I., Smidsrod, O., & Skjak-Braek, G. (2005).

Polysaccharides in the food industry: properties, production, and patents. Weinhein: John Wiley.

10. van de Ven, W. J. C., van’t Sant, K., Punt, I. G. M., Zwijnenburg, A., Kemperman, A. J. B., van der Meer, W. G. J., & Wessling, M. (2008). Hollow fiber dead-end ultrafiltration: Influence of ionic environment on filtration of alginates. Journal of Membrane Science, 308(1–2), 218-229. http://dx.doi.org/10.1016/j. memsci.2007.09.062.

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11. Carvalho, R. B., Borges, C. P., & Nóbrega, R. Formação de membranas planas celulósicas por espalhamento duplo para os processos de nanofiltração e osmose inversa. Polímeros: Ciência e Tecnologia, 11(2), 65-75. http://dx.doi.org/10.1590/ S0104-14282001000200008.

12. Shi, Y. Q., Wang, X. W., Chen, G. W., Golemme, G., Zhang, S. M., & Drioli, E. (1998). Preparation and characterization of high-performance dehydrating pervaporation alginate membranes.

Journal of Applied Polymer Science, 68(6), 959-968. http:// dx.doi.org/10.1002/(SICI)1097-4628(19980509)68:6<959::AIDAPP9>3.0.CO;2-G. 13. Dubois, M., Gilles, K. A., Hamilton, J. K., Rebers, F. A., & Smith, F. (1956). Colorimetric method for determination of sugars and related substances. Analytical Chemistry, 28(3), 350-356. http://dx.doi.org/10.1021/ac60111a017.

14. Hargrove, M. S., Wilkinson, A. J., & Olson, J. S. (1996). Structural factors governing hemin dissociation from metmyoglobin.

Biochemistry, 35(35), 11300-11309. http://dx.doi.org/10.1021/ bi960372d. PMid:8784184.

15. Ordway, G. A., & Garry, D. J. (2004). Myoglobin: an essential hemoprotein in striated muscle. The Journal of Experimental Biology, 207(Pt 20), 3441-3446. http://dx.doi.org/10.1242/ jeb.01172. PMid:15339940.

16. Marmo Moreira, L., Lima Poli, A., Costa-Filho, A. J., & Imasato, H. (2006). Pentacoordinate and hexacoordinate ferric hemes in acid medium: EPR, UV-Vis and CD studies of the giant extracellular hemoglobin of Glossoscolex paulistus.

Biophysical Chemistry, 124(1), 62-72. http://dx.doi.org/10.1016/j. bpc.2006.05.030. PMid:16814451.

17. Mahlicli, F. Y., Altinkaya, S. A., & Yurekli, Y. (2012). Preparation and characterization of polyacrylonitrile membranes modified with polyelectrolyte deposition for separating similar sized proteins. Journal of Membrane Science, 415-416(1), 383-390. http://dx.doi.org/10.1016/j.memsci.2012.05.028.

Received: June 26, 2014 Revised: Oct. 12, 2014 Accepted: Nov. 18, 2014

Polímeros , 25(3), 265-270, 2015

http://dx.doi.org/10.1590/0104-1428.1780

Wear and friction of composites of an epoxy with boron containing wastes Tayfun Uygunoğlu1, Witold Brostow2 and Ibrahim Gunes3 Civil Engineering Department, Faculty of Engineering, Afyon Kocatepe University, 03200, Afyonkarahisar, Turkey 2 Laboratory of Advanced Polymers & Optimized Materials – LAPOM, Department of Materials Science and Engineering, Center for Advanced Research and Technology – CART, University of North Texas, 76207, Denton, TX, USA 3 Department of Metallurgical and Materials Engineering, Faculty of Technology, Afyon Kocatepe University, 03200 Afyonkarahisar, Turkey 1

*uygunoglu@aku.edu.tr

Abstract Polymer surface coatings provide superior adhesion to substrates, some flexibility and corrosion resistance. On the other hand, 400,000 ton of boron wastes are generated each year. We have developed polymer composites based on epoxy resins containing up to 50 wt. % of boron wastes and determined their pin-on-disk dynamic friction, wear, Shore D hardness and surface roughness. The hardness and wear resistance increase with increasing boron waste concentration.

An equation, with parameters dependent on the load, relating wear rate to hardness is provided. Dynamic friction increases with increasing surface roughness, as represented by the equation. Further, dynamic friction is an increasing function of the wear rate. Micrographs of pure epoxy without fillers shows traces after pin-on-disk testing, with tears, breaks and cracks. For the composites, we observe simpler and relatively homogeneous surfaces.

Keywords: boron-containing waste, epoxy composites, abrasive wear, dynamic friction, shore hardness, roughness.

1. Introduction The largest boron deposits are found in Turkey with a worldwide share of 72% in terms of B2O3 content. During the obtaining of boron minerals such as tincalconite (Na2O·2B2O3·5H2O), ulexite (Na2O·2CaO·5B2O3·16H2O) and colemanite (2CaO·3B2O3·5H2O), about 400,000 tons of different types of boron wastes are formed and rejected in tailing dams per year[1]. A few studies are carried out on evaluation of boron wastes. Kavas et al.[1] investigated the production of artificial lightweight aggregates (LWA) by using four boron-containing wastes (BW), named as Sieve (SBW), Dewatering (DBW), Thickener (TBW) and Mixture (MBW) waste, from Kırka Boron plant in Turkey.

They reported that SBW and DBW boron-containing wastes combined with a clay mixture and quartz sand can be used for the manufacturing of LWA. Kurama et al.[2] used a dewatering sieve waste (TSW) of Etibor Kırka Borax company (Turkey) in order to develop an experimental terracotta floor tile body composition in combination with a feldspathic waste provided from a local sanitaryware plant and a ball clay. The results indicated a prospect for using the TSW as a raw material in mixtures with both clay and sanitaryware waste for the production of a terracotta floor tile body. The utilization of boron-containing clay wastes as cement additives was investigated by Özdemir and Öztürk[3].

It was observed that the first clay wastes may be used as cement additives up to 5% or 10%.

On the other hand, polymeric materials are noted for their versatility, high resistance to chemicals, outstanding adhesion to a variety of substrates, toughness, high electrical

Polímeros, 25(3), 271-276, 2015

resistance, durability at high and low temperatures, low shrinkage upon cure, flexibility, and the ease with which they can be poured or cast without forming bubbles[4-6].

Various kinds of polymers and polymer–matrix composites reinforced with metal particles have a wide range of industrial applications such as heaters, electrodes[7], composites with thermal durability at high temperature[8], etc. The inclusion of such particulate fillers into polymers for commercial applications is primarily aimed at the cost reduction and stiffness improvement.

The wear behavior of polymeric materials has drawn a considerable interest in recent years. Polymers and their composites are being increasingly used in a various applications where resistance to abrasive wear is important[9].

These range from its use as a material (in applications such as machinery parts and biomedical joint replacements) to its use as a glazing material where damage results in loss of optical properties. Polymers are ideal materials for bearing applications due to their general resistance to corrosion, galling and seizure, their tolerance to small misalignments and shock loading and their low coefficients of friction; as glazing materials, their low density and high toughness along with high transparency are desirable properties[10,11].

The acceptability of polymeric materials for abrasive wear conditions largely depends upon its mechanical load carrying capacity and the wear rate. The practical choice of polymeric materials is however not only determined by the mechanical and tribological properties, but also by the price, simplicity of production, processing and the practical limitations in the

271

S S S S S S S S S S S S S S S S S S S S

Uygunoğlu, T., Brostow, W., & Gunes, I. real application[12,13]. The performance of polymers sliding against hard and smooth counterfaces is determined by the transfer ability and buildup of a polymer film. Efficiency of materials in reducing friction and wear depends on the molecular polymer structure and counterface type. However, only few publications are available on the comparison of the tribological properties of composites under dry sliding and abrasive wear conditions[14-18].

In general, studies take into account to enhance the wear resistance of polymer materials[19-22]. Epoxy resins are the most commonly used thermoset plastic in polymer matrix composites which do not give off reaction products when they cure and thus have low cure shrinkage[4-6]. They also have good adhesion to other materials, good chemical and environmental resistance, good chemical properties and good insulating properties. We have investigated the influence of boron-containing wastes as a filler on wear and friction characteristics of epoxy composites.

2. Experimental 2.1. Materials and sample preparation The boron-containing wastes used in the study were provided by the Eti Holding Borax Plant in Kırka/Eskişehir, and were taken from the outlet of dewatering sieve of dissolution units. Its maximum particle size was 500 µm.

Chemical component of waste is presented in Table 1.

Commercially available Teknobond 300 epoxy resin along with hardener was used as matrix material in fabrication of different specimens. Epoxy resin has modulus of 3.42 GPa, and possess density of 1100 kg/m3. For processing, the mix ratio by weight was: 2 parts of the epoxy resin and 1 part of the amine based hardener. The required mixture of resin and hardener (see Table 2 below) were made by mixing them in a beaker by stirring the mixture by a rod, taking care that no air should be entrapped inside the solution. The composites were created at the room temperature.

The required ingredients of resin, hardener and boron waste were mixed thoroughly and the mixture so made was transferred to a mold cavity coated with a separator.

Steel moulds in size Ø = 50 mm were used for casting of the specimens. Curing was done at room temperature for approximately 24 h. After curing, the specimens were

de‑molded. Shore D hardness of the specimens was measured at 8 different locations at the same distance from the surface and averages calculated.

2.2. Tests on polymer composites Dynamic friction and abrasive wear were determined with a tribometer in a ball-on-disc configuration. WC + Co balls of 2.0 mm diameter from H.C. Starck Ceramics, GmbH, Munich, were used. Experiments were carried out under a dry friction condition at room temperature with applied loads of 5.0, 10.0 and 15.0 N and with the sliding speed of 0.2 m/s at a sliding distance of 125 m. Before and after each wear test, each sample and abrasion element was cleaned with alcohol. After the tests, the wear volumes of the samples were quantified by multiplying cross-sectional areas of wear by the width of the wear track obtained from the device Tribotechnic Rugosimeter, namely Wear rate = W = Worn volume / (Applied load x Sliding distance), mm3/Nm

(1)

Dynamic friction values as function of the sliding distance were obtained using dedicated software. Surface profiles of the wear tracks on the samples and surface roughness were measured by a Tribotechnic Rugosimeter.

All the tests were performed on the three specimens and the averages calculated.

3. Results and Discussion 3.1. Surface roughness Figure 1 shows the surface roughness values. In samples with high volume waste content, the roughness is lower than in the control pure epoxy specimens, except for 10% boron waste.

3.2. Wear and hardness Wear rates of our composites are displayed in Figure 2 for the three loads applied as a function of the boron content.

As seen in that Figure, wear of composites decreases when increasing the content of boron-containing waste material. The highest wear values are obtained for control series at each loading conditions. The range of wear rates varies from 17x10 –5 to 2x10 –5 mm3/Nm at 5.0 N, from

Table 1. Chemical content of boron wastes.

Oxide

B2O3

SiO2

Al2O3

CaO

MgO

K2O

Na2O

Fe2O3

Content, %

12.09

15.5

1.38

17.7

13.79

0.5

3.34

0.22

Loss of Ignition (LOI) 34.4

Table 2. Compositions and hardness of the composites.

Mixture code B0 B10 B20 B30 B40 B50

Epoxy resin*, kg/m3 100 90 80 70 60 50

Boron waste, kg/m3 10 20 30 40 50

Hardness (HD) 32 38 47 59 74 98

*Epoxy resin was used with the hardener in the 2:1 ratio.

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Polímeros , 25(3), 271-276, 2015

Wear and friction of composites of an epoxy with boron containing wastes 21x10 –5 to 3x10 –5 mm3/Nm at 10.0 N, from 27x10 –5 to 3.5x10 –5 mm3/Nm at 15.0 N, depending on waste content. In other words, there is approximately 7.7 times enhancement in wear resistance of polymer composites at 15 N caused by the waste presence, with higher waste concentrations producing larger effects. Thus, the addition of the waste material makes the epoxy material harder and also changes the character of the surface. Wear strength is higher at higher waste concentrations due to more homogeneous distribution of filler particles since they enhance the resistance to abrasion.

At the same time, when load values are considered, the wear rate for all specimens increases with load for all waste material ratios. The wear rate increases with an increase of load from 5 to 15 N due to increase of abrasion and friction on surface of polymer composite.

A relationship between the wear rate and hardness has been investigated and is shown in Figure 3. For each load, the wear resistance increases with an increase of the hardness. The respective equations are provided in inserts to Figure 3. The coefficient of determination R 2 is very close to 1.0 for each loading condition, the highest for 15.0 N (R 2 = corresponds to the perfect fit).

3.3. Friction Steady state dynamic friction values at different loads are reported in Table 3. The highest dynamic friction is seen in the control samples, namely 0.4. In the composite with 50% waste the friction is equal to 0.21, a reduction to almost one half. As expected, increasing the load increases dynamic friction.

An explanation of friction results can be provided in terms of ‘‘bumps’’ on predominantly polymer composite surface. The bumps model has been advanced before[23] and later confirmed in several cases[24,25]. Without the filler, the nominal contact area is equal to the real contact area.

Addition of the filler results in bumps formation; more filler means lower real contact area. Increasing the load makes the bumps slightly less effective in friction lowering.

We have also pursued in Figure 4 the dependence of the dynamic friction values on the surface roughness. There is a clear correlation between these quantities. R 2 values are here reasonably good, the best fit is seen for 15 N. In turn, we present the relation between dynamic friction and wear rate in Figure 5. A clear correlation emerges in Figure 5: = F 0.164 + 0.0103W (2)

Here F is the dynamic friction; W is the wear rate as before, calculated from Equation 1.

3.4. Microstructure

Figure 1. Surface roughness values of epoxy composites with boron waste.

We observed specimens after ball-on-disk tribometry at 15 N using an optical microscope and SEM; see Figures 6 and 7, respectively. Importantly, waste material particles are uniformly distributed in the epoxy matrix. In Figures 6a and 7a, the cracks and deformations can be clearly seen on the surface of the control (pure epoxy) specimen. The largest track width Table 3. Dynamic friction values under different conditions.

Figure 2. Wear rates of epoxies for various boron waste concentrations.

Figure 3. Relationship between hardness and wear rate of the composites in the inserts.

Polímeros, 25(3), 271-276, 2015

Mixture code B0 B10 B20 B30 B40 B50

5.0 N 0.37 0.28 0.26 0.21 0.19 0.16

10.0 N 0.39 0.33 0.30 0.27 0.22 0.19

15.0 N 0.40 0.36 0.32 0.30 0.24 0.21

Figure 4. Relationship between dynamic friction and surface roughness.

273

Uygunoğlu, T., Brostow, W., & Gunes, I. is also observed for pure epoxy, possibly due to small air bubbles in the mixtures during the hardening. Needless to say, the microscopy results confirm the findings reported above, particularly in Figure 2.

We have seen in Table 2 that more waste as the filler increases the hardness. The filler also increases the wear resistance, as seen in Figure 3. These facts are reflected in Figures 6 and 7. Figure 6 shows us how the pin moves across ‘borders’ between the phases; the arrows illustrate such movements.

Figure 5. Relationship between dynamic friction and wear rate.

Pure epoxy without fillers shows traces with tears, breaks and cracks. For the composites we see simpler and relatively homogeneous surfaces.

Figure 6. Microstructures of epoxy composites after tribometry for various boron waste concentrations: (a) for pure epoxy, (b) for 10% waste, (c) for 20% waste, (d) for 30% waste, (e) for 40% and (f) for 50% waste.

Figure 7. The wear SEMicrographs of epoxy composites after tribometry for various boron waste concentrations: (a) for pure epoxy, (b) for 10% waste, (c) for 20% waste, (d) for 30% waste, (e) for 40% and (f) for 50% waste.

274

Polímeros , 25(3), 271-276, 2015

Wear and friction of composites of an epoxy with boron containing wastes

4. Conclusions We report above wear and friction characteristics of epoxy based polymer composites as well as the role of boron-containing waste additives on properties. Several conclusions are derived: • The hardness values of samples are increased three times by using boron-containing waste material when compared to pure epoxy.

• The surface roughness values are similar to each other in samples with or without waste material content; the lowest roughness is seen for 50 wt. % waste.

• Dynamic friction decreases with increasing waste concentration while it increases slightly with increasing load.

• Wear resistance is higher in high waste material composites, apparently due to more homogeneous distribution of filler particles with high resistance to abrasion. The wear rate for the 50 wt. % boron waste composite is 7.7 times smaller than for the pure epoxy.

• A good was relationship between hardness versus wear rate; coefficient of friction versus roughness; and coefficient of friction versus wear rate is obtained, respectively, for epoxy based and waste material blended polymer composites.

• Microstructure observations show that the waste material particles are uniformly distributed in the epoxy matrix.

While the deformations such as tears, breaks and cracks are observed in pure epoxy specimens, there is only trace on the waste material containing samples after the tribometry tests.

5. Acknowledgements The author would like to thank National Boron Research Institute (BOREN) of Turkey, Ankara, for the financial support under the Project No. 2013.Ç0405.

6. References 1. Kavas, T., Christogerou, A., Pontikes, Y., & Angelopoulos, G. N. (2011). Valorisation of different types of boron-containing wastes for the production of lightweight aggregates. Journal of Hazardous Materials, 185(2-3), 1381-1389. http://dx.doi. org/10.1016/j.jhazmat.2010.10.059. PMid:21075514.

2. Kurama, S., Kara, A., & Kurama, H. (2006). The effect of boron waste in phase and microstructural development of a terracotta body during firing. Journal of the European Ceramic Society, 26(4-5), 755-760. http://dx.doi.org/10.1016/j. jeurceramsoc.2005.07.039.

3. Özdemir, M., & Öztürk, N. U. (2003). Utilization of clay wastes containing boron as cement additives. Cement and Concrete Research, 33(10), 1659-1661. http://dx.doi.org/10.1016/ S0008-8846(03)00138-8.

4. Brostow, W., Dutta, M., & Rusek, P. (2010). Modified epoxy coatings on mild steel: tribology and surface energy. European Polymer Journal, 46(11), 2181-2189. http://dx.doi.org/10.1016/j. eurpolymj.2010.08.006.

5. Bilyeu, B., Brostow, W., & Menard, K. P. (2001). Determination of volume changes during cure via void elimination and Polímeros, 25(3), 271-276, 2015

shrinkage of an epoxy prepreg using a quartz dilatometry cell.

Polimery, 46(11-12), 799-802. Retrieved from http://ichp.pl/ attach.php?id=2849 6. Bilyeu, B., Brostow, W., & Menard, K. P. (2001). Epoxy thermosets and their applications. III. Kinetic equations and models. Journal of Materials Education, 23(4-6), 189-204.

Retrieved from http://www.unt.edu/LAPOM/publications/ pdf%20articles/Lisa/epoxyJME3.pdf 7. Jang, B. Z. (1994). Advanced polymer composites: principles and applications. Ohio: Metals Park/ASM International.

8. Kim, J., Kang, P. H., & Nho, Y. C. (2004). Positive temperature coefficient behavior of polymer composites having a high melting temperature. Journal of Applied Polymer Science, 92(1), 394-401. http://dx.doi.org/10.1002/app.20064.

9. Harsha, A. P. (2011). An investigation on low stress abrasive wear characteristics of high performance engineering thermoplastic polymers. Wear, 271(5-6), 942-951. http://dx.doi.org/10.1016/j. wear.2011.03.019.

10. Shipway, P. H., & Ngao, N. K. (2003). Microscale abrasive wear of polymeric materials. Wear, 255(1-6), 742-750. http:// dx.doi.org/10.1016/S0043-1648(03)00106-6.

11. Cayer-Barrioz, J., Mazuyer, D., Kapsa, Ph., Chateauminois, A., & Robert, G. (2004). Abrasive wear micromechanisms of oriented polymers. Polymer, 45(8), 2729-2736. http://dx.doi. org/10.1016/j.polymer.2004.02.013.

12. Ginzburg, B. M., Tochil’nikov, D. G., Bakhareva, V. E., Anisimov, A. V., & Kireenko, O. F. (2006). Polymeric materials for water-lubricated plain bearings. Russian Journal of Applied Chemistry, 79(5), 695-706. http://dx.doi.org/10.1134/ S1070427206050016.

13. Gunes, I., Uygunoglu, T., Ergen, A., Kısıkcılar, T., & Aksoy, E. (2015). Investigation of wear behavior of borided DIN 20MoCr4 steel. El-Cezerî Journal of Science and Engineering, 2, 53-58. Retrieved from http://ecjse.com/index.php/ECJSE/ article/view/53 14. Samyn, P., Schoukens, G., Quintelier, J., & De Baets, P. (2006).

Friction, wear and material transfer of sintered polyimides sliding against various steel and diamond-like carbon coated surfaces. Tribology International, 39(6), 575-589. http://dx.doi. org/10.1016/j.triboint.2005.07.029.

15. Cirino, M., Friedrich, K., & Pipes, R. B. (1988). Evaluation of Polymer Composites for Sliding and Abrasive Wear Applications.

Composites, 19(5), 383-392. http://dx.doi.org/10.1016/00104361(88)90126-7.

16. Cirino, M., Pipes, R. B., & Friedrich, K. (1987). The abrasive wear behaviour of continuous ibre polymer composites.

Journal of Materials Science, 22(7), 2481-2492. http://dx.doi. org/10.1007/BF01082134.

17. Cirino, M., Friedrich, K., & Pipes, R. B. (1988). The effect of fiber orientation on the abrasive wear behavior of polymer composite materials. Wear, 121(2), 127-141. http://dx.doi. org/10.1016/0043-1648(88)90038-5.

18. Zhao, G., Hussainova, I., Antonov, M., Wang, Q., & Wang, T. (2013). Friction and wear of fiber reinforced polyimide composites. Wear, 301(1-2), 122-129. http://dx.doi.org/10.1016/j. wear.2012.12.019.

19. Rao, M., Hooke, C. J., Kukureka, S. N., Liao, P., & Chen, Y. K. (1998). The effect of PTFE on the friction and wear behaviour of polymers in rolling-sliding contact. Polymer Engineering and Science, 38(12), 1946-1958. http://dx.doi.org/10.1002/ pen.10364.

20. Pihtili, H. (2009). An Experimental Investigation of Wear of Glass Fibre-Epoxy Resin and Glass Fibre-Polyester Resin Composite Materials. European Polymer Journal, 45(1), 149154. http://dx.doi.org/10.1016/j.eurpolymj.2008.10.006.

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Uygunoğlu, T., Brostow, W., & Gunes, I. 21. Svancarek, P., Lendvayova, S., Galusek, D., Hnatko, M., Vavra, I., & Wang, X. (2011). Abrasive wear resistance of SiO2-doped polycrystalline alumina. Wear, 271(5-6), 760-769. http://dx.doi.org/10.1016/j.wear.2011.03.016.

22. Brostow, W., Kovacevic, V., Vrsaljko, D., & Whitworth, J. (2010). Tribology of polymer and polymer-based composites.

Journal of Materials Education, 32(5-6), 273-290.

23. Brostow, W., Kumar, P., Vrsaljko, D., & Whitworth, J. (2011).

Optimization of tribological and mechanical properties of nanocomposites of polyurethane/poly(vinyl acetate)/CaCO3.

Journal of Nanoscience and Nanotechnology, 11(5), 3922-3928. http://dx.doi.org/10.1166/jnn.2011.3849. PMid:21780387.

24. Olea-Mejía, O., Brostow, W., Escobar-Alarcón, L., & ViguerasSantiago, E. (2012). Tribological properties of polymer

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25. Brostow, W., Datashvili, T., & Geodakyan, J. (2012). Tribological properties of ethylene–propylene–diene rubber + polypropylene + thermal-shock-resistant ceramic composites. Polymer International, 61(9), 1362-1370. http://dx.doi.org/10.1002/ pi.4282.

Received: May 23, 2014 Revised: Sept. 09, 2014 Accepted: Nov. 18, 2014

Polímeros , 25(3), 271-276, 2015

http://dx.doi.org/10.1590/0104-1428.1621

Analysis of equations of state for polymers Erlí José Padilha Júnior1, Rafael de Pelegrini Soares1 and Nilo Sérgio Medeiros Cardozo1 Departamento de Engenharia Química – DEQUI, Universidade do Rio Grande do Sul – UFRGS, CEP 90040-040, Porto Alegre, RS, Brasil

1

*nilo@enq.ufrgs.br

Abstract In the literature there are several studies comparing the accuracy of various models in describing the PvT behavior of polymers. However, most of these studies do not provide information about the quality of the estimated parameters or the sensitivity of the prediction of thermodynamic properties to the parameters of the equations. Furthermore, there are few studies exploring the prediction of thermal expansion and compression coefficients. Based on these observations, the objective of this study is to deepen the analysis of Tait, HH (Hartmann-Haque), MCM (modified cell model) and SHT (simplified hole theory) equations of state in predicting the PvT behavior of polymers, for both molten and solid states. The results showed that all equations of state provide an adequate description of the PvT behavior in the molten state, with low standard deviations in the estimation of parameters, adequate sensitivity of their parameters and plausible prediction of specific volume, thermal expansion and isothermal compression coefficients. In the solid state the Tait equation exhibited similar performance to the molten state, while HH showed satisfactory results for amorphous polymers and difficulty in adjusting the PvT curve for semicrystalline polymers.

Keywords: equation of state, PvT behavior, polymer.

1. Introduction The study of the thermodynamic behavior of polymers is essential to analyze the physical transformations that occur during processing, e.g. injection molding or extrusion, and to predict the properties of the final products. Polymers in melt state or in solution can be represented correctly by an equation of state (EoS) because these can be considered equilibrium states for polymers. Conversely, the solid state is at least one quasi-equilibrium state, because the properties depend on the conditions of solidification, as the cooling rate and pressure, which difficults their description by means of EoS[1].

Numerous equations of state have been developed to describe the PvT (pressure-volume-temperature) behavior of polymers. In literature, there are several studies comparing the fitting accuracy of various models for PvT data of polymers[1-7].

These studies are focused mainly in the analysis of the fitting accuracy of the specific volume in the molten state, employing usually the method of least squares[2,3,6,8-20] for the parameter estimation step. They showed that the theoretical equations based on cell-and‑hole models and the Tait and the Hartmann-Haque (HH) empirical equations are those which provide more accurate fitting of the experimental data.

On the other hand, little information is available in the literature on at least two important aspects that are essential for the use of the referred equations in process simulation: quality of the estimated parameters as a function of the model employed and fitting accuracy of the models for thermal expansion and isothermal compression coefficients.

The isobaric thermal expansion coefficient ( β ) and isothermal compressibility ( κ ) are defined by: 1  ∂ν  (1) β=   ν  ∂T  P

Polímeros, 25(3), 277-288, 2015

1  ∂ν  κ = −   (2) ν  ∂P T

where the negative sign indicates the volume decrease with pressure increase[21].

These coefficients are important for the simulation of polymer processing operations, because they are present in the governing equation of energy conservation. However, to the best of our knowledge, the only studies on their prediction from EoS are those of Utracki[22,23], in which the prediction of thermal expansivity and compressibility by hole models is analyzed.

Based on these observations, the objective of this study is to deepen the analysis of Tait, HH, MCM (modified cell model) and SHT (simplified hole theory) equations of state in prediction of PvT behavior of polymers, in both molten and solid physical states. The EoS were analyzed with respect to: (i) quality in the estimation of its parameters by the method of least squares, (ii) sensitivity of their predictions to each of its parameters, (iii) quality of the prediction of the specific volume, and (iv) quality of the prediction of isobaric thermal expansion coefficient and isothermal compressibility.

2. Equations of State Analyzed 2.1 Tait equation of state This equation is purely empirical, and was originally proposed for water. Presently, through various modifications, it is applied to a wide variety of substances, being possibly one of the equations of state most used to model the PvT behavior of polymers[3]. For some authors, it is not a true equation of state, but an isothermal compressibility model

277

S S S S S S S S S S S S S S S S S S S S

Padilha, E. J., Jr., Soares, R. P., & Cardozo, N. S. M. (i.e., a volume-pressure relationship). Tait equation can be written as[3]:   P  + v T , P ) (3) v (T , P )= vo (T ) 1 − Cln 1 +  B (T )   t (    

where for polymers in the molten state, i.e., above the liquid-solid transition temperature: vo = b1m + b2m (T − b5 ) (4a)

= P

P v T (9) = ; v = ; T B0 s v0 s T0 s

where B0 , v0 and T0 are the characteristic parameters.

T0 and v0 are defined as temperature and specific volume, respectively, extrapolated to zero pressure, while B0 is identified as the isothermal bulk modulus extrapolated to zero temperature and pressure.

2.3 MCM equation of state B (T )= b3m exp  −b4 m (T − b5 )  (4b)

The modified cell model equation of state was developed

by Dee and Walsh[2], starting from the formalism presented

by Prigogine et al.[25]. In the cell model, the compressibility νt (T , P ) = 0 (4c) and thermal expansion of the structure are explained only by changes in the cell volume. Dee and Walsh[2] introduced and for polymers in solid state, i.e., below the liquid-solid a numerical factor that scales the hard-core cell volume in transition temperature the free volume term, disconnecting the theory from the specific geometry. This factor, q , was found to be constant vo =+ b1s b2 s (T − b5 ) (5a) for numerous polymers and equal to about 1.07. MCM EoS can be written as: B (T )= b3s exp  −b4 s (T − b5 )  (5b)  Pv v1/3 2  1.2045 1.011  = 1/3 −  – 4  (10)  T v – 0.8909q T  v 2 v  νt= (T , P ) b7exp {b8 (T − b5 )  − ( b9 P )} (5c) The reduced parameters P , v and T are defined as: The liquid-solid transition temperature, which is the v  T  P (11) = P = ; v = ;T glass transition temperature for amorphous polymers and * * P v T* the melting or crystallization temperature for semicrystalline polymers, can be calculated by where P* , v* and T * are characteristic parameters.

Tt ( P= ) b5 + b6 P (6)

2.4 SHT equation of state

The hole theory introduces empty cells in the cell model[26], In these equations, v is the specific volume of the based on the concept that the thermal expansion of liquid is polymeric material; the coefficient C is a constant equal mainly due to holes ( h ), i.e., the empty cells, while volume to 0.0894; vo is the specific volume at zero pressure; νt changes of the cells are also allowed. Zhong et al.[27] simplified is the specific volume corresponding to crystalline phase; the hole theory through the use of an exponential function B is the sensitivity to pressure of material; b1 at b9 are to the fraction of occupied cells. SHT EoS is derived as: parameters of model, obtained by fitting of PvT diagram.

The parameters b1m to b4m and b1s to b4s describe the   ( yv )1/3 Pv 2 y  1.1394 dependence on pressure and temperature in the molten and   (12) = + –1 .5317  T ( yv )1/3 – 0.9165 y T ( yv )2  ( yv )2 the solid state, respectively; b5 and b6 are parameters that   describe the change of transition temperature with pressure; b7 to b9 are particular parameters of semicrystalline polymers where P , v and T are reduced parameters defined by that describe the form of the state transition[24].

Equation 11, y is the fraction of occupied cells, being defined by:

2.2 HH equation of state

Hartmann and Haque[10] developed an empirical equation of state combining the thermal pressure function of Pastine and Warfield, the zero-pressure isobar presented by Somcynsky and Simha, and the empirical dependence of volume with the thermal pressure. HH EoS describes the PvT behavior of polymers in the molten and solid states. It is given by:   5 T 3/ 2 − lnv (7) Pv =

where the dimensionless variables P , v and T for molten polymers are defined as: P v T = P = ; v = ; T (8) B0m v0m T0m and for solid polymers as: 278



y = 1 − e−0.52/ T (13)

3. Methodology The experimental PvT data used in this work were taken from the literature[7,21,28-32] as shown in Table 1. For each polymer, the available data were subdivided in two sets: one used for parameter estimation (DATA1) and other for validation (DATA2). The construction of these subsets was based on random selection of points.

Differently from specific volume data, the isobaric thermal expansion and isothermal compression coefficients are hard to obtain experimentally. Thus, the EoS were only qualitatively analyzed with relation to the prediction of these Polímeros , 25(3), 277-288, 2015

Polﾃｭmeros, 25(3), 277-288, 2015

468

0.1

33

28

33

28

Initial temperature of analysis (K)

Initial pressure of analysis (MPa)

Number of experimental points in solid state used for parameter estimation

Number of experimental points in molten state used for parameter estimation

Number of experimental points in solid state used for the analysis of the prediction accuracy

Number of experimental points in molten state used for the analysis of the prediction accuracy

0.001

32

24

30

26

0.1

313

313-603

0.1-200

29

PC

0.0004

-

-

20

36

0.1

353

353-423

0.1-180

30

PMMA

0.0004

26

30

26

30

0.1

357

357-453

0.1-180

30

PCHMA

Amorphous

0.0004

-

-

22

29

0.1

285

285-355

0.1-180

30

PnBMA

0.0004

-

-

26

25

0.1

302

302-470

0.1-180

31

PoMS

0.001

24

28

25

28

0.1

313

313-573

0.1-200

29

iPP

0.002

-

-

29

41

0.1

313

313-493

0.1-200

32

LPE

a

Data obtained by the method of confining fluid in isothermal mode, except for polystyrene, for which the experiments were conducted in isobaric mod.

2 ( ﾏテxp )

0.0003

280-468

Temperature range (K)

Experimental data variance

0.1-200

28

PS

Pressure range (MPa)

Reference

Polymer

Table 1. Relevant information about the experimental PvT data useda.

0.002

-

-

43

27

0.1

313

313-493

0.1-200

7

PEO

0.0003

17

35

17

35

0.1

364

364-586

0.1-190

21

PA6

Semicrystalline

0.002

-

-

21

56

0.1

313

313-493

0.1-200

32

PLA

0.002

19

15

20

16

0.1

313

313-493

0.1-200

32

PBS-B

Analysis of equations of state for polymers

279

Padilha, E. J., Jr., Soares, R. P., & Cardozo, N. S. M. coefficients, taking as basis of comparison their theoretically expected behavior. The estimation of parameters was conducted by the least squares method, using the lsqnonlin function already implemented in MatLab software, with the following objective function ( FObj ): n

2

FObj = ∑ ( vˆi − vi ) (14) i =1

where ( vˆi − vi ) is the residual between predicted ( vˆi ) and experimental ( vi ) values and n is the number of points considered.

The parameters of the equations of state were estimated simultaneously, as suggested by Hartmann and Haque[10].

For the Tait EoS, firstly, b5 and b6 were estimated from data of transition temperature at different pressures. The melt parameters ( b1m , b2m , b3m and b4m ) and solid parameters ( b1s , b2s , b3s , b4s , b7 , b8 and b9 ) were estimated separately from data corresponding to the respective states[33]. Likewise, for HH EoS, the experimental data were divided into two states: molten and solid state. The parameters of the MCM and SHT EoS were estimated only with molten state data.

The quality of the estimated parameters for each model was analyzed in terms of their covariance matrix ( Vα ), evaluated using a routine developed in MatLab, according the following expression[34]:

( )

= Vα H α−1Gα σ2y GαT H α−1

T

(15)

where σ2y is the experimental data variance, H α is the Hessian matrix of FObj , and Gα is the matrix that represents the derivative of the gradient of FObj with relation to the experimental values vi .

The normalized parameter sensitivity matrix was used to evaluate the sensitivity of the predictions of the considered EoS to parameter variations. The coefficients of this matrix are given by: Sij* =

∂vˆi  a j    (16) ∂a j  vˆi 

where a j is the parameter of the equation analyzed. These coefficients were calculated using a MatLab function, according to the following central-difference approximation: Sij* =

(

)

(

)

vi α j + h − vi α j − h  α j  2h  vi

 (17)  

with h equal to 10-4. In this way, the parameter sensitivity of the specific volume predictions was analyzed in a wide temperature range, at different pressures, for polycarbonate and linear polyethylene.

To appraise the fit and prediction quality of each model, the mean relative deviation (MRD) and the regression 2 coefficient ( R ) were calculated: MRD =

100 n vi − vˆi (18) ∑ n i =1 vi

R2 = 1 − 280

2

∑in=1 ( vi − vˆi ) (19) 2 ∑in=1 ( vi − vi )

where vi is the arithmetic mean of experimental specific volumes.

Additionally, F-tests were performed to assess the suitability of the considered models. The value of F0 used in the comparison with the critical value of F ( Fc = F distribution value corresponding to 95% confidence) was defined as 2 ) and the experimental the ratio between the model ( σeos 2 ) variances, according to the following expression: ( σexp 2

= F0

2 σeos = 2 σexp

∑in ( vi − vˆi ) (20) n − np 2 σexp

where np is the number of estimated parameters.

The experimental variances are shown in Table 1.

4. Results and Discussion 4.1 Fitting and prediction of specific volume As mentioned previously, the comparison of the EoS under analysis with relation to the fitting of specific volume data has already been extensively studied by other authors[2,3,6,8-20]. Therefore, in the present work, the analysis corresponding to the fitting stage will be focused only in the quality of the estimated parameters, aspect for which there is little information available in the literature.

The values of parameters estimated from data set DATA1 and their respective standard deviations are presented in Tables 2 and 3, for the amorphous and semicrystalline polymers, respectively. The low values of the standard deviations indicate that the fitting were adequate. The parameters of equations of state showed a higher standard deviation for semicrystalline polymers, in both physical states. The fitting of Tait and HH EoS were better in molten state, both to amorphous and semicrystalline polymers. In general, Tait equation exhibited the lowest values of standard deviations (below 2%) except for the parameter b6 in the cases of the polymers iPP and PLA, for which the linear dependence between transition temperature and pressure described by Equation 6 is not obeyed. For all equations of state, the greatest deviations were found for the parameters related to pressure, what can be explained by the wide range of pressure analyzed.

For the evaluation of the prediction capability of the studied models, calculations of specific volume for the conditions corresponding to each experimental data of data set DATA2 were performed using the values of parameters estimated with data set DATA1 (Tables 2 and 3). The values of F0 , MRD and R 2 obtained are presented in Table 4, together with the respective values of Fc used in the F-test to 95% confidence. As can be seen in Table 4, all the four EoS provided predictions not significantly different from the experimental data ( F0 < Fc ), showing their adequacy in the prediction of the PvT behavior of the considered polymers. It is observed that Tait equation exhibited the lowest relative deviation module mean and the highest regression coefficient in most cases. However, the other equations of state studied also presented satisfactory results, with values​​ close to those obtained with Tait EoS. The only exception was in the prediction of specific volume of semicrystalline Polímeros , 25(3), 277-288, 2015

Analysis of equations of state for polymers Table 2. Estimated parameters (ai) and percentage standard deviation (σai) for amorphous polymers.

EoS Tait

HH

MCM

SHT

Parameter b1m(cm3/g)

PS

ai

PC

σai (%)

0.9767

0.0015

ai

σai (%)

0.8590

0.0025

PMMA

ai

σai (%)

0.8603

0.0024

PCHMA

ai

σai (%)

0.9318

0.0017

PnBMA

ai

σai (%)

0.9469

0.0014

PoMS

ai

Mean

σai (%) σai (%) 0.0013

0.0018

b2m(cm3/gK) 0.000506 0.0461

0.000553 0.0392

0.000511 0.1042

0.000596 0.0477

0.000628 0.0532

1.0046

0.000547 0.0492

0.0566

b3m (MPa)

154.65

0.0672

151.39

0.1062

277.56

0.1512

161.98

0.0629

192.89

0.0998

184.57

0.0882

0.0959

b4m (1/K)

0.0030

0.2749

0.0034

0.2218

0.0075

0.6696

0.0052

0.1845

0.0047

0.4290

0.0062

0.2388

0.3364

b1s (cm3/g)

0.9748

0.0015

0.8575

0.0035

0.8625

0.0007

0.9303

0.0014

0.9479

0.0011

1.0045

0.0018

0.0017

b2s (cm3/gK) 0.000213 0.1116

0.000192 0.2435

0.000275 0.1665

0.000248 0.1118

0.000376 0.3084

0.000232 0.0996

0.1736

b3s (MPa)

275.87

0.0645

249.21

0.1005

257.75

0.0303

255.61

0.0486

223.53

0.0373

276.20

0.0649

0.0577

b4s (1/K)

0.0018

0.6080

0.0021

0.7384

0.0045

0.3877

0.0038

0.3082

0.0036

0.8858

0.00067 1.2603

0.6981

b5 (K)

370.86

0.0026

417.06

0.0059

375.30

0.0034

383.84

0.0030

296.66

0.0039

401.21

0.0028

0.0036

b6 (K/MPa) 0.3924

0.0211

0.2687

0.0885

0.3485

0.0692

0.2665

0.0441

0.2225

0.0528

0.3361

0.0350

0.0518

B0m (MPa)

2953.3

1.3157

3470.2

2.6453

5238.2

5.0184

3115.0

1.3573

3654.1

2.5917

3154.5

2.1292

2.5096

v0m (cm3/g) 0.8753

0.6226

0.7413

1.3425

0.7570

0.7592

0.8192

0.7058

0.8431

0.6388

0.8917

0.6503

0.7865

T0m (K)

1602.1

0.2884

1471.8

0.6255

1464.9

0.7126

1471.8

0.3327

1256.7

0.4040

1631.2

0.3724

0.4559

B0s (MPa)

4278.5

3.6067

3858.2

7.8331

3867.7

3.2444

4172.8

3.7264

3499.0

3.3230

4251.8

4.6353

4.3948

v0s (cm3/g)

0.9218

0.5102

0.8107

1.4375

0.8158

0.5568

0.8785

0.6056

0.8891

0.7577

0.9399

0.6413

0.7515

T0s (K)

2552.0

0.4872

2914.7

1.3898

2592.6

0.4136

2656.6

0.5196

1863.6

0.4846

2453.5

0.7396

0.6724

P*(MPa)

532.04

0.4534

707.66

0.9018

1017.5

2.1312

594.83

0.8393

656.45

0.6733

616.17

0.9196

0.9864

v*(cm3/g)

0.8690

0.0822

0.7357

0.1447

0.7463

0.2813

0.8101

0.1353

0.8392

0.1177

0.8745

0.1612

0.1537

T*(K)

6933.9

0.3985

6441.5

0.5280

6178.8

1.1607

6298.7

0.5389

5427.0

0.5621

6704.6

0.6911

0.6466

P*(MPa)

420.57

1.0068

675.78

2.1852

887.13

6.4028

518.74

2.1551

506.85

1.3689

577.23

4.3201

2.9065

v*(cm3/g)

0.9242

0.1769

0.7620

0.3835

0.7798

0.8736

0.8491

0.3381

0.8937

0.2329

0.9032

0.7175

0.4538

T*(K)

4009.2

0.8431

3353.9

1.3560

3311.9

3.4399

3425.4

1.3222

3168.3

1.1022

3429.9

2.7175

1.7968

Table 3. Estimated parameters (ai) and percentage standard deviation (σai ) for semicrystalline polymers.

EoS Tait

Parameter

iPP

LPE

PEO

PA6

PLA

PBS-B

ai

σai (%)

ai

σai (%)

ai

σai (%)

ai

σai (%)

ai

σai (%)

ai

b1m (cm3/g)

0.0022

1.2301

0.0018

1.0002

0.0030

0.8875

0.0043

0.8837

b3m (MPa)

Mean

σai (%) σai (%)

1.3082

0.0021

1.2532

0.0035

0.0028

b2m (cm3/gK) 0.0010

0.0378

0.000957 0.0509

0.000932 0.0304

0.000659 0.0873

0.000712 0.1455

0.000622 0.0788

0.0718

66.84

0.0438

93.68

0.0769

113.12

0.0577

132.00

0.0860

100.68

0.1041

152.55

0.1278

0.0827

b4m (1/K)

0.0048

0.1116

0.0046

0.2521

0.0044

0.1411

0.0029

0.5355

0.0047

0.5810

0.0040

0.4383

0.3433

b1s (cm3/g)

1.1804

0.0033

1.0669

0.0030

1.1666

0.0975

0.9600

0.0020

0.8553

0.0025

0.8361

0.0109

0.0199

b3s (MPa)

110.82

b2s (cm3/gK) 0.000517 0.0743

0.000453 0.1063

0.000780 1.1831

0.000489 0.0396

0.000372 0.0695

0.000525 0.2850

0.2930

0.0765

228.35

0.0825

173.63

1.3530

125.96

0.0496

148.00

0.0902

198.59

0.3480

0.3333

b4s (1/K)

0.0064

0.1556

0.0023

0.5972

0.0029

3.0337

0.0078

0.0830

0.0064

0.1984

0.0057

1.1048

0.8621

b5 (K)

452.86

0.0055

405.47

0.0061

365.09

0.0068

501.95

0.0020

440.44

0.0056

385.88

0.0064

0.0054

b6 (K/MPa)

0.0057

4.1917

0.2107

0.1129

0.2043

0.1164

0.0835

0.1127

0.0048

4.9477

0.1254

0.1897

1.6119

b7 (cm3/g)

0.3644

0.2083

0.4848

0.1569

0.0565

1.9711

0.0406

0.0709

0.0327

0.1187

0.0133

2.7994

0.8876

b8 (1/K)

0.1429

0.1600

0.2893

0.1153

0.0210

1.4194

0.0890

0.0718

0.4684

0.2048

0.1785

1.4186

0.5650

b9 (1/MPa)

0.1133

0.2824

0.0762

0.1072

0.0074

0.5776

0.0029

0.2204

0.0081

0.2998

0.0111

1.3604

0.4746

B0m (MPa)

1877.2

0.9060

2514.9

1.5167

2581.2

1.3639

4330.2

2.6157

3235.6

1.6826

3648.7

3.1454

1.8717

v0m (cm3/g)

1.0864

1.5161

1.0496

1.4684

1.0605

1.1865

0.8137

1.5780

0.7196

1.8242

0.7573

1.8053

1.5631

T0m (K)

1370.7

0.3162

1273.9

0.4684

1277.8

0.4502

1437.2

0.7388

1242.7

0.6546

1327.6

0.8649

0.5822

B0s (MPa)

2689.0

3.2847

5508.9

5.8402

5242.2

2.0723

5634.4

1.7457

4112.4

4.1033

6393.9

9.2255

4.3786

v0s (cm3/g)

1.0241

1.4413

0.9064

1.0336

0.9097

1.6756

0.7828

0.5857

0.7426

1.0285

0.6945

2.0437

1.3014

T0s (K)

1574.8

0.9417

1284.6

0.9761

830.57

0.5881

1396.3

0.3584

1576.0

0.9510

1170.2

1.4734

0.8815

MCM P*(MPa)

406.30

1.4660

497.88

1.7795

516.67

0.9554

882.19

1.3540

668.30

3.7141

717.11

1.6693

1.8231

v*(cm3/g)

1.0656

0.1957

1.0435

0.2740

1.0492

0.1417

0.8110

0.1867

0.7141

0.5219

0.7510

0.2474

0.2612

T*(K)

5825.1

0.5207

5601.7

0.8363

5521.4

0.4740

6392.2

0.4969

5457.8

1.3499

5774.1

0.8529

0.7551

P*(MPa)

445.72

3.5869

509.45

4.7145

509.45

2.4290

922.41

3.4977

714.20

9.5570

678.21

4.2874

4.6788

v*(cm3/g)

1.0880

0.5217

1.0730

0.7374

1.0836

0.3770

0.8285

0.4958

0.7288

1.3734

0.7783

0.6555

0.6935

T*(K)

2929.4

1.3930

2864.6

2.2278

2855.9

1.2422

3209.4

1.3034

2737.2

3.5220

3016.0

2.2253

1.9856

HH

SHT

Polímeros, 25(3), 277-288, 2015

281

282

8.38×10–4 4.84×10–4 6.11×10–4 0.0017 0.0129 1.66×10–4 -

-

0.9994

F0

0.9997 0.9994 0.9998 0.9996 0.9978 0.9998 -

R

2

Tait

0.9992 0.9990 0.9984 0.9988 0.9980 0.9996 -

R

2

HH

0.0998 0.9988

MRD (%) 0.0581 0.1077 0.1136 0.1600 0.1010 0.0583 -

1.81×10–9 7.27×10–4 4.03×10–9 0.0043 0.0101 2.48×10–4 -

F0

F0

-

0.9997 6.38×10–4 0.9976 0.0018 0.9992 0.0020 0.9991 0.0031 0.9992 0.0040 0.9995 2.75×10–4 -

R 2

MCM

0.0867 0.9991

MRD (%) 0.0323 0.1333 0.0808 0.1320 0.0804 0.0612 -

Molten EoS F0

-

0.9995 0.0013 0.9988 9.08×10–4 0.9995 0.0014 0.9990 0.0035 0.9987 0.0067 0.9994 3.40×10–4 -

R 2

SHT

0.0924 0.9992

MRD (%) 0.0520 0.1141 0.0612 0.1478 0.1059 0.0733 -

1.4792 1.4465 1.4984 1.5200 1.6253 1.5891 -

Fca

-

MRD (%) 0.0201 0.0557 0.0639 0.4300 0.1964 0.0890 0.0466 0.2385 F0

-

-

0.9998 1.78×10–4 0.9982 1.97×10–4 0.9976 0.0020 0.9249 0.0933 0.9920 0.0297 0.9978 5.40×10–4 0.9985 0.9716 -

R 2

Tait

-

MRD (%) 0.0294 0.1395 0.1085 0.9429 0.6487 0.1741 0.0925 0.5886

Solid EoS F0

-

-

0.9994 4.22×10–4 0.9877 0.0012 0.9937 0.0049 0.8659 0.1532 0.9316 0.2388 0.9899 0.0021 0.9936 0.9291 -

R2

HH

-

1.4393 1.5200 1.4620 1.4792 1.4259 1.6689 -

Fcb

a

Degrees of freedom used: 28, 32, 26, 24, 17 and 19 for PS, PC, PCHMA, iPP, PA6 and PBS-B, respectively; bDegrees of freedom used: 33, 24, 30, 28, 35 and 15 for PS, PC, PCHMA, iPP, PA6 and PBS-B, respectively.

MRD (%) PS 0.0394 PC 0.0723 PCHMA 0.0364 IPP 0.1028 PA6 0.1070 PBS-B 0.0444 Amorphous Mean Semicrystalline Mean Mean 0.0671

Polymer

Table 4. Statistical results in specific volume prediction by EoS.

Padilha, E. J., Jr., Soares, R. P., & Cardozo, N. S. M.

Polímeros , 25(3), 277-288, 2015

Analysis of equations of state for polymers polymers in solid state, where a higher difference between Tait and HH equations occurred.

Figure 1 shows the residual plots for each equation of state. It is possible to observed that the predominately random nature of the errors distributions for both molten and solid state data, with the predictions of the HH EoS

for molten iPP as only relevant exception. These results support the statement of good suitability of the EoS tested.

As example of the general behavior described in the previous paragraph, Figure 2 shows the variation of MRD with the temperature for an amorphous polymer, PC, and with the pressure and temperature for a semicrystalline one, iPP. It can be seen that the HH EoS presented a high relative

Figure 1. Residual plots of specific volume predictions for all EoS tested: (a) Molten and (b) Solid state polymers.

PolĂ­meros, 25(3), 277-288, 2015

283

Padilha, E. J., Jr., Soares, R. P., & Cardozo, N. S. M.

Figure 2. MRD (%) in specific volume prediction: (a) PC in 10 MPa, (b) PC in 200 MPa, (c) iPP in 443.8 K and (d) iPP in 20 MPa.

deviation in solid state, especially at low pressures near the transition temperature (Figures 2c and 2d).

4.2 Sensitivity analysis The sensitivity of the prediction of specific volume with relation to each parameter of HH, MCM and SHT EoS is shown in Figure 3. The sensitivity of parameters showed similar behavior for all these EoS, both to amorphous and semicrystalline polymers. Besides the highest sensitivity corresponds to the volume related parameters, the sensitivity to each parameter was nearly constant in the whole ranges of pressure and temperature analyzed.

For the Tait EoS, the behavior of the sensitivity to the parameters was somewhat different, as shown in Figure 4.

The sensitivity to the parameters b1 and b2 varied continuous and complementarily with the increase of the temperature (Figures 4b and 4c), with increase of the sensitivity to b1 and decrease of the sensitivity to b2 . Moreover, it is perceived that there was an abrupt change in the sensitivity of the parameters of Tait EoS near the transition temperature of linear polyethylene in the solid state. The sensitivities of parameters b1s and of term vt ( b7 , b8 and b9 ) of Tait equation were modified near the transition region. It is 284

found that there is a correlation between the sensitivities to the parameters b1s and b7 , both related to the specific volume. All parameters of the vt term ( b7 , b8 and b9 ) reveal sensitivity close to the transition regions, stating that they are within linked in modeling this region. Thus, this transition region has fundamental importance for the estimation of the parameters of the Tait EoS.

4.3 Isobaric thermal expansion and isothermal compression coefficients prediction analysis Figure 5 shows the isobaric thermal expansivity calculated from the PC, PoMS, iPP and PLA by equations of state. It is observed that all the EoS predict similar values of this coefficient and are in qualitative agreement with the theory in the sense that the thermal expansion coefficient of a polymer melt is always greater than that of the corresponding amorphous and semicrystalline solid[35,36]. However, Tait EoS, unlike the other equations, predicts a reduction of this coefficient with the increase of the temperature, contrarily to theoretical expectations[35], revealing a limitation of the model. Moreover, in the case of semicrystalline polymers, Tait equation of state presented an abrupt increase in thermal expansion coefficient in the Polímeros , 25(3), 277-288, 2015

Analysis of equations of state for polymers crystalline transition region. This occurs because Tait EoS describes satisfactorily sudden change in specific volume due to destruction/growth of crystallites, which does not happen with the HH equation. The theoretical equations of state, MCM and SHT displayed the same curve shape.

Figure 6 shows the isothermal compressibility predicted from the PC, PoMS, iPP and PLA by equations of state.

It appears that the equations predict values near. The curves exposed by EoS are consistent with the theory[35,36], coefficient gradually increases with temperature and decreases with pressure. Again, as in the case of predicting thermal expansivity, Tait equation of state exhibited an abrupt reduction of isothermal compressibility in the crystalline transition region.

Figure 3. Normalized sensitivity of specific volume relative to the parameters of: (a) HH EoS to PC in 0.1 MPa, (b) HH EoS to LPE in 200 MPa, (c) MCM EoS to PC in 0.1 MPa, (d) MCM EoS to LPE in 200 MPa, (e) SHT EoS to PC in 200 MPa and (f) SHT EoS to LPE in 0.1 MPa. Polímeros, 25(3), 277-288, 2015

285

Padilha, E. J., Jr., Soares, R. P., & Cardozo, N. S. M.

Figure 4. Normalized sensitivity of specific volume of polymers relative to parameters of Tait equation of state: (a) PC in 0.1 MPa, (b) LPE in 0.1 MPa, (c) LPE in 200 MPa and (d) LPE in 405.9 K.

Figure 5. Isobaric thermal expansion coefficient predicted by EoS: (a) PC, (b) PoMS, (c) iPP and (d) PLA.

286

Polímeros , 25(3), 277-288, 2015

Analysis of equations of state for polymers

Figure 6. Isothermal compressibility predicted by EoS: (a) PC, (b) PoMS, (c) iPP and (d) PLA.

5. Conclusions

6. References

The Tait, HH, MCM and SHT equations of state were evaluated in prediction the PvT behavior of polymers, for both molten and solid physical states.

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2. Dee, G. T., & Walsh, D. J. (1988). A modified cell model equation of state for polymer liquids. Macromolecules, 21(3), 815-817. http://dx.doi.org/10.1021/ma00181a044.

3. Rodgers, P. A. (1993). Pressure–volume–temperature relationships for polymeric liquids: A review of equations of state and their characteristic parameters for 56 polymers. Journal of Applied Polymer Science, 48(6), 1061-1080. http://dx.doi.org/10.1002/ app.1993.070480613.

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5. Wang, W., Liu, X., Zhong, C., Twu, C. H., & Coon, J. E. (1997). Simplified hole theory equation of state for liquid polymers and solvents and their solutions. Industrial & Engineering Chemistry Research, 36(6), 2390-2398. http:// dx.doi.org/10.1021/ie9604132.

6. Chiew, Y. C., Ting, S. K. H., & Leong, K. K. (2000). A perturbed Lennard–Jones chain equation of state for polymer liquids. Fluid Phase Equilibria, 168(1), 19-29. http://dx.doi. org/10.1016/S0378-3812(99)00329-5.

7. Sato, Y., Hashiguchi, H., Inohara, K., Takishima, S., & Masuoka, H. (2007). PVT properties of polyethylene copolymer melts. Fluid Phase Equilibria, 257(2), 124-130. http://dx.doi. org/10.1016/j.fluid.2007.01.013.

In the analysis of the PvT behavior of melt polymers, all equations of state studied showed adequate fitting of specific volume data, with light advantage of the Tait equation.

No significant differences among them were observed in terms of quality of the estimated parameters, sensitivity of the predictions to the parameters, and of prediction of the thermal expansion and compression coefficients. Then, all EoS studied are appropriate in modeling the molten state.

In the analysis of the PvT behavior of solid polymers, Tait and HH equations exhibited differences in sensitivity analysis and specific volume prediction, justified mainly because HH EoS does not describe correctly the crystalline transition. The parameter estimation in both equations was adequate, with low values ​​of standard deviations. Thus, the Tait equation of state is the most appropriate for modeling solid polymers, except for the prediction of the isobaric thermal expansion coefficient, property for which the values predicted with this equation are not in qualitative agreement with theoretical expectations.

Based on these results, the Tait equation of state can be indicated as the most appropriate for modeling the PvT behavior during processing of polymers.

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9. Sanchez, I. C., & Lacombe, R. H. (1978). Statistical thermodynamics of polymer solutions. Macromolecules, 11(6), 1145-1156. http://dx.doi.org/10.1021/ma60066a017.

10. Hartmann, B., & Haque, M. A. (1985). Equation of state for polymer solids. Journal of Applied Physics, 58(8), 2831-2836. http://dx.doi.org/10.1063/1.335881.

11. Ougizawa, T., Dee, G. T., & Walsh, D. J. (1989). PVT properties and equations of state of polystyrene: molecular weight dependence of the characteristic parameters in equationof-state theories. Polymer, 30(9), 1675-1679. http://dx.doi. org/10.1016/0032-3861(89)90329-7.

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Received: Feb. 01, 2014 Revised: Oct. 22, 2014 Accepted: Dec. 04, 2014

Polímeros , 25(3), 277-288, 2015

http://dx.doi.org/10.1590/0104-1428.1840

Acacia bark residues as filler in polypropylene composites Ticiane Taflick1,2, Élida Gonçalves Maich3, Laís Dias Ferreira1, Clara Isméria Damiani Bica1, Silvia Rosane Santos Rodrigues2 and Sônia Marlí Bohrz Nachtigall1 Federal University of Rio Grande do Sul – UFRGS, CEP 91501-970, Porto Alegre, RS, Brazil 2 Sul-rio-grandense Federal Institute for Education, Science and Technology – IFSUL, CEP 93216-120, Sapucaia do Sul, RS, Brazil 3 La Salle University, CEP 92010-000, Canoas, RS, Brazil

1

*sonia.nachtigall@ufrgs.br

Abstract Large amounts of acacia bark residues are produced each day after tannin extraction with hot water, being generally burned. This by-product was chemically characterized and used as filler in polypropylene (PP) composites, considering different particle sizes and concentrations. The materials produced by melt blending had their mechanical and thermal properties evaluated. It was verified that, even containing a significant amount of extractable compounds, the acacia bark particles can produce PP composites with higher impact properties, higher crystallization temperature and higher degradation temperature in comparison to the polymer matrix.

Keywords: composites, natural fibers, impact properties, tensile properties, thermal degradation.

1. Introduction Polymers filled with lignocellulosic materials have been employed since the 80’s in the form of wood plastic composites (WPCs). Natural fibers obtained from plants are environmentally friendly alternatives as reinforcement for synthetic polymers mainly due to their relative high strength and stiffness, low cost, low density, low CO2 emission, biodegradability and renewable character[1]. They are non‑abrasive during processing, allowing the improvement of stiffness and strength in thermoplastics[2,3]. Although WPCs have been widely studied[4-7], a survey of the literature reveals that the use of bark particles as thermoplastic fillers has not been extensively evaluated.

The bark is the outer part of the tree trunks and branches.

It shows different chemical composition and is less fibrous than the woody parts of the trees[8, 9]. As a residue from industries, the bark is mostly used for thermal energy production. Its use as an alternative raw material for WPCs might result in better utilization of such natural resource.

The black acacia (Acacia mearnsii De Wild) is a tree species from Australia which grows fast and with a plentiful of uses (bark and wood). It is one of the most important cultivated trees in the state of Rio Grande do Sul, Brazil.

The hot water extract obtained from black acacia bark is rich in tannins, which are polyphenols largely used for leather tanning, water treatment, and for pharmaceutical and chemical products filtering[9]. Nowadays, the bark residue obtained after tannin extraction is an environmental problem and it has been generally burned. Thus, the use of black acacia bark as filler in polymers can be an alternative to an environmental damage caused by the tannin industry, leading to both economic and environmental benefits.

Among the commodities used in the manufacture of WPCs, polypropylene (PP) is very important because of its low cost, specific low weight and good mechanical

Polímeros, 25(3), 289-295, 2015

properties for non-structural applications[10]. It has been shown that both bark and wood flour significantly increase the mechanical strength of thermoplastics, irrespectively of the botanical species[11]. However, the effect of bark particles is less pronounced. Safdari et al.[11] attributed this behavior to the presence of fines, to the low aspect ratio of bark particles and to the lower intrinsic fiber strength of bark fibers compared to wood fibers.

The present study agrees with the global trend regarding the decrease in consumption of raw materials from non‑renewable sources and with the efforts to reduce global warming. So, we have studied the mechanical and thermal properties of PP composites prepared with black acacia bark after tannin extraction. Polypropylene functionalized with maleic anhydride (MAPP) was used as a compatibilizer to improve the adhesion between the matrix and the filler[5,12-14].

2. Experimental 2.1. Materials The black acacia bark, after tannin extraction with hot water followed by compression, was supplied by SETA S/A ‑ Extrativa Tanino de Acácia localized in Estância Velha, Brazil.

Sulfuric acid, hexane and ethanol p.a. from Fmaia, Brazil, were used as received. Isotactic PP (MFI = 3.5 g/10 min; d = 0.905 g/cm3) was supplied by Braskem S.A.. PP modified with maleic anhydride (MAPP, 1% maleic anhydride) Polybond 3200 from Uniroyal Chemical was used as coupling agent.

2.2. Filler conditioning The black acacia bark was dried in a first step for 12 h, at 60 °C, and in a second step at 70 °C, in a dryer oven, under low pressure, until constant weight. The material was triturated using a cryogenic rotatory mill, and separated

289

S S S S S S S S S S S S S S S S S S S S

Taflick, T., Maich, E. G., Ferreira, L. D., Bica, C. I. D., Rodrigues, S. R. S., & Nachtigall, S. M. B. using ASTM standard mesh sieves. Three different ranges of particle size were used: short (106-212 μm), medium (212-425 μm) and large (425-600 μm).

2.3. Preparation of the composites The composites were prepared in a melt mixer Haake, Rheomix 600p, at 50 rpm, for 8 min, at 180 °C, according to the compositions shown in Table 1. All components were introduced at the same time into the mixer.

Films were prepared using a hydraulic press Carver, Monarch series, 3710C model, by pre-heating samples at 180ºC (5 min) followed by pressing at 2500 kgf (5 min).

Injected specimens for tests were obtained using a Thermo Scientific MiniJet II, by injecting small composite pieces at 180 °C. The injection pressure used was 250 bar, with 5s injection and mould temperature at 40oC.

Calorimetric analyses were performed using a TA Instruments Q20, under inert atmosphere, at 10 °C/min heating rate, from room temperature to 180 °C. The samples were kept at this temperature for 3 min and then were cooled down to 25 °C at 60 °C/min and heated again to 200 °C at 10 °C/min. The reported thermal data were taken during the cooling and second heating cycles. The degree of crystallinity (Xc, %) was determined from the second melting enthalpy considering the melting enthalpy of 100% crystalline PP as 209J/g[16].

Thermogravimetric analyses were performed using a TA Instrument Q50, at 20°C/min heating rate, from 50 to 800 °C, under inert atmosphere.

3. Results and Discussion 3.1. Filler characterization

2.4. Bark acacia characterization The as-received acacia bark particles were imaged using a conventional camera. SEM images of the ground acacia bark particles were done using a JEOL microscope JSM 6060, operating at 20 kV. The particles were attached to an aluminum stub and sputtered with gold to eliminate electron charging effects.

3.1.1. Morphology The raw acacia bark contained a mixture of large flat particles and thin fibers, as shown in Figure 1a. Because of this fact, some particles show high aspect ratio (> 50) while others have aspect ratio around one.

The content of extractives in acacia bark was determined through Soxhlet extraction with hexane, ethanol and water, in this sequence, according to TAPPI T 204 (adapted).

The lignin content was determined by the method described in TAPPI T 222, by measuring the insoluble fraction in sulfuric acid. The ash content was determined by calcination in a muffle furnace Provecto Analítica MFLO1000, according to TAPPI T 211: the material was gradually heated to 525 °C, and then left at that temperature until constant weight.

The holocellulose content (cellulose + hemicellulose) was determined by difference. Ground samples, in triplicate, were dried out until constant weight to determine the moisture content.

X-ray diffraction analysis was performed using a Siemens D-500 diffractometer, using an incident X-ray of CuKα with wavelength of 1.54 Å. The Segal method[15] was used to calculate the crystallinity of cellulose present in the sample according to Equation 1, where XC is the sample crystallinity. I200 is the height of the 200 peak, which represents both crystalline and amorphous cellulose and IAM is the lowest height between the 200 and 110 peaks, which represents only the amorphous fraction.

XC = (I200 – IAM/I200) x 100%

(1)

2.5. Characterization of the composites Samples for Izod impact tests were prepared by injection moulding according to ASTM D256, in rectangular samples with 63.5 (±2) mm length, 12.7 (±0.2) mm width and 3.18 mm thick. The notched was formed at 45° with a depth of 2.45 mm.

Tensile tests were performed according to ASTM D638 in a Universal Testing Machine EMIC DL 10000, using injected samples with 62.6 mm of length.

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Figure 1. (a) Image of the as-received acacia bark particles; (b) SEM micrograph of ground acacia bark particles (sample LF).

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Acacia bark residues as filler in polypropylene composites Figure 1b (Sample LF) shows that although a significant amount of fines could be produced during grinding, they were not observed in the samples. This image shows many fibrous structures together with a minor fraction of irregular‑shaped small particles. Besides, the surface of the particles seemed to be very smooth probably due to the hot water treatment during tannin extraction, which could remove small components originally adhered to the material.

3.1.2. Chemical composition of acacia bark Studies in the literature have shown that some kinds of barks present lower cellulose content, higher lignin content and higher hot and cold water extractives than their woods[17,18].

They may still have high inorganic content, mainly due to the presence of silica, which results in high ash amount[11].

The contents of extractives, holocellulose, lignin, ash and residual moisture in acacia bark particles were determined.

The results are presented in Table 2.

As can be seen in Table 2, the percentage of extractives in acacia bark was relatively high (26.3%). However, in comparison to poplar bark (51%)[11] it was low. As the acacia bark used had been previously submitted to hot water treatment during tannin extraction, part of the soluble material had been certainly removed. From the total extractives, 16.8% were soluble in organic solvent and 9.5% were soluble in water. The remaining water-soluble compounds could be now extracted due to the smaller size of the particles used.

The low content of holocellulose (42.5%) in black acacia bark reflects a low content of cellulose, when compared to common woods (57-63%)[19]. However, these results come up well with findings in the literature for acacia bark[17]. The high lignin content (20.5%) can potentially contribute to an excellent rigidity when compared to other bark fibers[20]. The low level of ash found for acacia bark in our study (2.1%), indicates that this material shows low inorganic content.

3.1.3. X-ray diffraction Among the acacia bark components, only cellulose is crystalline. In this study, X-ray diffraction was used to evaluate the degree of crystallinity of cellulose in black acacia bark. Figure 2 shows that the major crystalline peaks of cellulose in acacia bark correspond to those found in materials exhibiting Cellulose I structure, where the peak around 2θ = 22° represents the (2 0 0) crystallographic plane while the peak around 2θ = 15° is assigned to the (1 0 1) crystallographic plane[21].

The crystallinity index calculated by Segal formula for black acacia bark was 60%. Lower values of crystallinity have been found in the literature for other bark samples, such as 49.3% for Eucalyptus grandis, 43.4% for Pinus elliottii[19] and 46% for Prosopis Juliflora barks[20]. This result suggests that acacia bark particles can be better used as filler in polymer composites since higher crystallinity is expected to improve mechanical properties.

3.2. Preparation and characterization of composites Usually, commercial WPCs formulations contain particles that range in size from 180 to 1,700 micrometers[22].

Stark and Rowlands[23] observed that smaller wood fibers performed significantly better than coarser fibers in both flexural and tensile strength and they attributed this trend to larger stress concentrations along the naturally weak interface of the larger wood particles. However, it was also found that systems containing larger fibers could show better performance than others containing smaller fibers if an interfacial agent was present[22]. In the present study, PP composites containing different sizes of acacia bark particles were prepared through melt processing, as shown in Table 1.

Considering the addition of PP-g-MA as a coupling agent, the effect of different amounts of the larger-sized filler was investigated. A composition containing only the polymer and the coupling agent PPMA was also prepared for comparison.

3.2.1. Mechanical properties

Table 1. Formulations of the composites.

Identification PP* SF-10 MF-10 LF-10 LF-20 LF-30

PP

MAPP

(%, wt)

(%, wt)

98 88 88 88 78 68

2 2 2 2 2 2

Filler content (%, wt) 0 10 10 10 20 30

Filler size (μm)

It was found that the impact strength of the composites containing particles with size from 106 to 425 μm exceeded that of the polymer matrix (Sample PP*, Figure 3), as a consequence of additional mechanisms of energy absorption

106-212 212-425 425-600 425-600 425-600

Table 2. Chemical composition of black acacia bark.

Black acacia bark Extractives (%)

26.3

Holocellulose (%)

42.5

Lignin (%)

20.5

Ash (%)

2.1

Moisture (%)

8.6

Polímeros, 25(3), 289-295, 2015

Figure 2. X-ray diffractogram of black acacia bark.

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Taflick, T., Maich, E. G., Ferreira, L. D., Bica, C. I. D., Rodrigues, S. R. S., & Nachtigall, S. M. B. during fracture. Probably, the small filler particles behaved as toughening centers, blocking crack propagation under impact and generating this interesting behaviour. However, the use of particles in the range of 426-600 μm decreased the impact strength to a value below that of the polymer matrix. Similar results were found by Saini et al.[17], which encountered lower impact strength for larger acacia bark particles in PVC/acacia bark flour composites. According to Figure 3, the impact strength showed no dependence on the amount of filler used but it was solely related to the size

Figure 3. Effect of acacia bark particles on the Izod impact strength.

of the filler particles. Larger particles provide more stress along their larger interface with the polymer, allowing the impact fracture to propagate more easily. In conclusion, PP composites prepared with the smaller acacia bark particles were more attractive for impact properties.

On the other hand, the tensile strength of all composites containing 10% of filler was kept similar to the polymer matrix, showing no dependence on the size of particles (Figure 4a). This indicates a good balance of filler dispersion and adhesion[24]. The addition of 2 wt% coupling agent was efficient in these systems. A distinct behavior was found for higher amounts of the large-sized particles (20 and 30%), which showed somewhat lower tensile resistance. This is related to the presence of the filler itself and to the lower MAPP/filler ratio, since the MAPP content was kept constant in all compositions while the filler content was increased.

In a previous study it was verified that, for PP composites containing 30% of several lignocellulosic fibers, the tensile strength increased when the concentration of MAPP was increased up to 10%[13].

A significant decrease of the Young modulus and elongation at break was found with the addition of the fillers in comparison to the base matrix (PP*), according to Figures 4b and 4c. This behavior is partially due to the presence of the soluble molecules in the bark composition,

Figure 4. Tensile properties: (a) tensile strength; (b) Young’s modulus; (c) elongation at break.

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Acacia bark residues as filler in polypropylene composites represented by the solvent extractable fraction. Soluble materials decrease the ratio of structural components in the filler. Moreover, large amounts of extractives could cause a decrease in the polarity of the filler surface and a decrease in wettability, limiting the PPMA performance[25].

Saini et al. 17] found similar results in PVC/acacia bark composites having acacia bark particle sizes in the range 100-150 μm. Considering elongation at break, it was observed that the compositions containing small and medium-size particles behaved better. According to the literature[1], short and tiny fibers (average particle size 0.24-0.35 mm) should be preferred in polymer composites, since they provide higher specific surface area and the fibers are distributed more homogeneously compared to composites with long fibers. In fact, these materials do not show properties for use in high performance applications, but they can be used in specific applications where the mechanical strength is not so important, such as in decoration and other popular goods.

hydrophobic in nature. According to Yang et al.[26] the peak at 382 °C can be attributed to the maximum degradation rate of cellulose. The shoulder at the left of this peak is attributed to hemicellulose. Lignin decomposition starts at relatively low temperature (~170 °C), with low mass loss, and it happens on a wide range of temperature extending approximately to 600 °C[21].

For the polymer matrix it is seen that the degradation happens in one step with the maximum rate occurring near 450 °C. This maximum can be related to the breaking of C-C bonds in PP chains[27]. The presence of MAPP did not show any additional important degradation event. Curves in Figure 5 clearly indicate that PP increased its thermal stability after compounding. An increase of 42 °C in the highest peak of the weight loss derivative curve for the composite containing 30% of large size fillers was observed Table 3. Thermal properties of the composites.

3.2.2. Thermal properties Table 3 presents the thermal data obtained by DSC: crystallization temperature (TC), melting temperature (Tm), melting enthalpy (∆Hm) and crystallinity (XC).

PP melting temperature did not change in the composites, indicating that the crystalline structure of the polymer remained unaltered. However, the addition of filler increased polymer TC and XC, probably due to an effect of nucleation. High crystallization temperature and crystallinity are interesting properties for thermoplastic processors since they reduce the production time, saving costs. The crystalline fraction of PP in all composites was similar.

Representative TG and DTG curves are presented in Figure 5. Table 4 shows the temperatures of maximum degradation rates (Tmax) and the percentage of residue remaining at 800 °C in the samples.

A close look at the curves in Figure 5 shows that the filler starts to degrade at temperatures lower than the polymer matrix. The weight loss of acacia bark below 100 °C is explained by water loss and this event was not observed in composites, indicating they are mainly

Sample PP* SF-10 MF-10 LF-10 LF-20 LF-30 (1) (3)

Tm(1)

TC(2)

XC(3)

∆Hm

(°C) 166.5 166.6 166.2 166.2 166.9 166.7

(°C) 103.5 106.9 106.0 105.4 106.2 109.5

(%) 38 46 46 45 44 46

(J/g) 80 86 87 84 74 68

Maximum of the melting peak. (2) Onset of the crystallization peak.

Degree of crystallinity.

Table 4. Thermal data obtained from TGA (under nitrogen).

Sample Acacia bark PP* SF-10 MF-10 LF-10 LF-20 LF-30

Tmax

Residue at 800 °C

(°C) 382 453 464 455 459 470 495

(%) 22.40 0.06 1.60 1.70 1.60 1.90 3.70

Figure 5. TG and DTG curves of acacia bark, PP* and composite LF-30.

Polímeros, 25(3), 289-295, 2015

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Taflick, T., Maich, E. G., Ferreira, L. D., Bica, C. I. D., Rodrigues, S. R. S., & Nachtigall, S. M. B. (Table 4). We can attribute these results to the restriction of the polymer chains slippage in the presence of the fillers, which makes the mobility of the active species during degradation more difficult.

High residual mass was found for the acacia bark at the end of the analysis (22.40%). This is probably related to an incomplete degradation of the filler components at 800 °C, under nitrogen. In accordance to this value, the residual mass of the composites increased with the content of the filler. However, the mass of residues in the composites was not as high as expected considering the filler amount, meaning that the smaller dimensions of the particles in the composites favored their degradation.

4. Conclusions The results showed that PP composites prepared with black acacia bark particles after tannin extraction may have interesting properties, such as higher impact resistance, higher crystallization temperature and higher thermal stability than the matrix, depending on the amount and size of the filler used. Composites prepared with 10% acacia bark particles with sizes ranging from 106 to 425 µm showed better properties than those prepared with larger particles (425-600 µm) and higher amount (up to 30%). As a bark raw material, these fillers contain high amount of extractable compounds that negatively affect the rigidity and the plasticity of the composites. However, the tensile strength was kept unaltered for the best compositions. Due to this fact, they can be suitable for use in non-structural applications, such as in decoration, sound isolation and household goods, with the advantages of reducing the environmental impact of using synthetic polymers and adding value to a waste product formed after tannin extraction from black acacia bark.

5. Acknowledgements The authors thank to SETA S.A. for black acacia bark donation, to FAPERGS for scholarship and grants, to Federal University of Rio Grande do Sul, Sul-Rio‑Grandense Institute for Education, Science and Technology and La Salle University.

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14. Becker, D., Kleinschmidt, A. C., Balzer, P. S., & Soldi, V. (2011). Influência da sequência de mistura do PP-MA nas propriedades dos compósitos de PP e fibra de bananeira.

Polímeros: Ciência e Tecnologia, 21(1), 7-12. http://dx.doi. org/10.1590/S0104-14282011005000012.

15. Thygesen, A., Oddershede, J., Lilholt, H., Thomsen, A. B., & Stahl, K. (2005). On the determination of crystallinity and cellulose content in plant fibers. Cellulose, 12(6), 563-576. http://dx.doi.org/10.1007/s10570-005-9001-8.

16. Ayrilmis, N., & Kaymakci, A. (2013). Fast growing biomass as reinforcing filler in thermoplastic composites: Paulownia elongate wood. Industrial Crops and Products, 43, 457-464. http://dx.doi.org/10.1016/j.indcrop.2012.07.050.

17. Saini, G., Bharwaj, R., Choudhary, V., & Narula, A. K. (2010).

Poly(vinyl chloride)-Acacia bark flour composite: effect of particle size and filler content on mechanical, thermal, and morphological characteristics. Journal of Applied Polymer Science, 117(3), 1309-1318. http://dx.doi.org/10.1002/app.29987.

18. Yemele, M. C. N., Koubaa, A., Cloutier, A., Soulounganga, P., & Wolcott, M. (2010). Effect of bark fiber content and size on the mechanical properties of bark/HDPE composites.

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19. Poletto, M., Zattera, A. J., Forte, M. M. C., & Santana, R. M. C. (2012). Thermal decomposition of wood: influence of wood components and cellulose crystallite size. Bioresource Technology, 109, 148-153. http://dx.doi.org/10.1016/j. biortech.2011.11.122. PMid:22306076.

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Acacia bark residues as filler in polypropylene composites cellulosic fiber from Prosopis juliflora bark. Carbohydrate Polymers, 92(2), 1928-1933. http://dx.doi.org/10.1016/j. carbpol.2012.11.064. PMid:23399239.

21. Popescu, M.-C., Popescu, C.-M., Lisa, G., & Sakata, Y. (2011). Evaluation of morphological and chemical aspects of different wood species by spectroscopy and thermal methods.

Journal of Molecular Structure, 988(1-3), 65-72. http://dx.doi. org/10.1016/j.molstruc.2010.12.004.

22. Ghallager, L. W., & McDonald, A. G. (2013). The effect of micron sized wood fibers in wood plastics composites.

Maderas, Ciencia y Tecnología, 15(3), 357-374. http://dx.doi. org/10.4067/S078-221X2013005000028.

23. Stark, N. M., & Rowlands, R. E. (2003). Effects of wood fibers characteristics on mechanical properties of wood/polypropylene composites. Wood and Fiber Science, 35(2), 167-174.

24. Părpăriţă, E., Darie, R. N., Popescu, C. M., Uddin, M. D., & Vasile, C. (2014). Structure-morphology-mechanical properties relationship of some polypropylene/lignocellulosic composites.

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Materials & Design, 56, 763-772. http://dx.doi.org/10.1016/j. matdes.2013.12.033.

25. Najafi, S. K., Kiaefar, A., & Tajvidi, M. (2008). Effect of bark flour content on the hygroscopic characteristics of wood– polypropylene composites. Journal of Applied Polymer Science, 110(5), 3116-3120. http://dx.doi.org/10.1002/app.28852.

26. Yang, H., Yan, R., Chen, H., Zheng, C., Lee, D. H., & Liang, D. T. (2006). In-depth investigation of biomass pyrolysis based on three major components: hemicelluloses, cellulose and lignin. Energy & Fuels, 20(1), 388-393. http://dx.doi. org/10.1021/ef0580117.

27. Shebani, A. N., Van Reenen, A. J., & Meincken, M. (2009). The effect of wood species on the mechanical and thermal properties of wood–LLDPE composites. Journal of Composite Materials, 43(11), 1305-1318. http://dx.doi.org/10.1177/0021998308104548.

Received: July 11, 2014 Revised: Nov. 20, 2014 Accepted: Dec. 18, 2014

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http://dx.doi.org/10.1590/0104-1428.1961

S S S S S S S S S S S S S S S S S S S S

Melting and crystallization of poly(3-hydroxybutyrate): effect of heating/cooling rates on phase transformation Renate Maria Ramos Wellen1, Marcelo Silveira Rabello2, Inaldo Cesar Araujo Júnior3, Guilhermino José Macedo Fechine4 and Eduardo Luis Canedo2,5 Departamento de Engenharia de Materiais, Universidade Federal de João Pessoa – UFPB, CEP 58051-900, João Pessoa, PB, Brazil 2 Departamento de Engenharia de Materiais, Universidade Federal de Campina Grande – UFCG, CEP 58429-140, Campina Grande, PB, Brazil 3 Departamento de Engenharia Química, Universidade Federal de Pernambuco – UFPE, CEP 50740-521, Recife, PE, Brazil 4 Centro de Pesquisa sobre Grafeno e Nano-Materiais – Mackgraphe, Universidade Presbiteriana Mackenzie, CEP 01302-904, São Paulo, SP, Brazil 5 Instituto de Tecnologia de Pernambuco – ITEP, CEP 50740-540, Recife, PE, Brazil 1

*wellen.renate@gmail.com

Abstract We studied the crystallization and melting phenomena of poly (3- hydroxybutyrate) (PHB), a biodegradable and biocompatible semi-crystalline thermoplastic, obtained from renewable resources. Its high crystallinity motivated several studies on crystallization and melting behavior, and also on ways to increase the amorphous polymer fraction.

The effect of heating and cooling rates on the crystallization and melting of commercial PHB was investigated by differential scanning calorimetry. Several rates, ranging from 2.5 to 20 °C min–1, were used to study the phase changes during heating/cooling/reheating cycles. The results showed that PHB partially crystallizes from the melt during the cooling cycle and partially cold crystallizes on reheating, and that the relative amount of polymer crystallizing in each stage strongly depends on the cooling rate. The melt and cold crystallization temperatures, as well as the rates of phase change, depend strongly on the cooling and heating rates.

Keywords: PHB, DSC, crystallization, melting, kinetics.

1. Introduction Consumer products based on polymers are produced in large quantities to fulfill the needs of modern society.

Automotive, shipbuilding, textiles, electronic devices, food packing, healthcare, etc, are examples of industrial areas that use polymeric resins derived from natural or synthetic sources on the manufacture of products with specific properties. Most applications use conventional commodity and engineering thermoplastics. However, polymers produced from natural resources and biodegradable plastics are gaining increasing interest, not only in biomedical applications but also in packing and consumer products[1-3]. Among these types of polymers, Poly (3-hydroxybutyrate) (PHB) is a very promising material, a semi-crystalline thermoplastic polymer that is biodegradable and biocompatible, and is obtained from renewable resources (mainly sugar cane) by biotechnological processes of low environmental impact[4-7].

Ongoing research, intended on establishing PHB’s properties and optimize its processing conditions, are reported in the literature with an expectation of extending the range of its application. One of the major problems with PHB is its high crystallinity that allied with a relatively high glass transition temperature, results in a very fragile material.

This fact is an important limitation to the practical use of PHB. Material properties of semi-crystalline polymers are controlled by molecular and supra-molecular structures that

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are frequently determined by the crystallization mechanisms.

Consequently, the study of crystallization and melting behavior is critical to understand and control material properties and the processing required to obtain them[8-15].

An experimental program was developed in order to understand the effect of thermal cycles on the phase transitions of the polymer. In this work the crystallization and melting phenomena in PHB was investigated by differential scanning calorimeter (DSC), applying thermal cycles of heating/cooling/reheating at rates ranging from 2.5 °C/min to 20 °C/min. DSC tests revealed that PHB may be crystallized from molten state, phenomenon known as melt crystallization, as well from the rubbery amorphous solid state, phenomenon known as cold crystallization.

The characteristic phase change temperatures and temperature intervals, the rate of phase change, and the amount of each phase involved in the change were determined in terms of the cooling and heating rates.

2. Experimental 2.1. Materials The poly(3-hydoxybutyrate) (PHB) polymer, actually a random copolymer with approximately 4% 3- hydoxyvalerate units, was supplied by PHB Industrial SA (Brazil). Thermal

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Melting and crystallization of poly(3-hydroxybutyrate): effect of heating/cooling rates on phase transformation transition temperatures of PHB are presented in Table 1.

Melt crystallization (TC) and melting peak (TMP) temperatures were estimated during cooling and reheating, respectively, according to ASTM D-3418; values reported were obtained at cooling/heating rate of 10oC/min.

The molecular mass distribution was obtained with a Viscotek HT-GPC Module 350A GPC at 40°C, with a refractive index detector. The material was dissolved in chloroform and the filtered solution was injected into the equipment. Solvent flow rate was 1 mL/min and the columns were calibrated with narrow molecular weight polystyrenes.

The molecular mass distribution curve of PHB is shown in Figure 1. From these data the number-average molar mass MN = 52 kg/mol and the polydispersity index MW/MN = 2.66 were estimated.

2.2. Methods 2.2.1. Differential scanning calorimetry (DSC) measurements Thermal analysis was performed in a Shimadzu DSC‑60 differential scanning calorimeter, under a nitrogen flow of 50 mL/min to minimize oxidative degradation to which the PHB is prone[8,16]. Samples of approximately 5 mg were wrapped with aluminum foil to minimize the effect of the polymer thermal conductivity, which may cause differential broadening and shifting of peak positions[17].

A new specimen was used for each run. A blank curve was obtained for each heating/cooling/reheating stage to ensure that no contamination of the instrument had taken place. A thermal cycle in four stages was used: (1) heating from room temperature (approximately 30°C) to 195°C (first heating stage); (2) at this temperature the samples were held a constant temperature for 3 min to eliminate any residual crystallinity and erase previous thermal history; (3) after which the melt was cooled to −10°C (cooling stage), and then (4) reheated to 195°C (second heating or reheating stage). Tests were run at constant rates of heating and cooling of 2.5, 5, 7.5, 10, 15, and 20°C/min. Figure 2 shows typical DSC output with the indicated cycles. Four thermal events were identified in most DSC tests, namely, melting during the first heating stage (F1), melt crystallization during cooling (C1), cold crystallization during reheating (C2), and, finally, a second melting event (F2).

For each thermal event, the starting and end points of departure from the underlying baseline were visually established in a plot of energy flow (J) versus time (t).

The fractional crystallization (or melting) x for the event was computed as a function of time by integration: = x(t )

1 t ∫ J (t ′) − J 0 (t ′) dt ′ E0 t1

The rate of phase change (crystallization or melting) c is simply: J (t ) − J 0 (t ) x c(t ) ≡ dx =

dt

E0

(3)

from which the peak (maximum) and average crystallization rates may be computed. The fractional crystallization/melting x and the rate of crystallization/fusion c may be expressed as functions of temperature (T), knowing the linear relationship between time and temperature during the event: T= T1 + φ(t − t1 )

(4)

where T1 is the sample temperature at the starting point t1, and φ is the (constant) rate of heating or cooling during the event. The specific latent heat of crystallization or melting (or enthalpy, since the phase change occurs at constant pressure) is computed from E0, polymer fraction wP, and the sample mass mS: E ∆H = o w p ms

(5)

Figure 1. Molecular weight distribution curve of PHB used in this work.

(1)

where t1 and t2 were the initial and final times, J0 is the virtual baseline during the event (straight in the present case), and E0 is the total latent heat of phase change: = E0 ∫tt2 J (t ) − J 0 (t ) dt 1

(2)

Table 1. Thermal transition temperatures of PHB. Polymer PHB

TG (°C) 2

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TC (°C) 56.5

TMP (°C) 169

Figure 2. Typical DSC output for PHB heating/cooling/reheating at 5°C/min (exothermic peaks up), showing the phase change events: first melting (F1), melt crystallization (C1), cold crystallization (C2), and second melting (F2).

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Wellen, R. M. R., Rabello, M. S., Araujo, I. C., Jr., Fechine, G. J. M., & Canedo, E. L. The mass crystallinity change X during the event is estimated taking into account the heat of fusion of PHB 100% crystalline ∆H°m: ∆H ∆X = o ∆H m

(6)

The value ∆H°M = 146 J/g at the equilibrium melting temperature T °M = 185°C was reported in the literature[18].

2.2.2. Heating and cooling rates In heat flow DSC instruments the nominal heating and cooling rates are, at best, aproximations to the rates of change of the reference temperature. True heating and cooling rates were computed from the sample temperature output of the DSC as: Ti +1 − Ti −1

φ(ti ) = 2∆t

(7)

where φ is the rate at time ti; Ti−1, Ti+1 are the temperatures at times ti − ∆t and ti + ∆t, and ∆t = 1 s is time interval between measurements. True heating/cooling rates computed for the temperature intervals of the second fusion (heating) and melt crystallization (cooling) and presented in Table 2, along with their 95% confidence intervals. The deviations (σ ) correspond to the RMS average point-to-point variation.

Table 2. True heating/cooling rates (in °C/min). ϕnom 2.5 5.0 7.5 10 15 20

On Cooling ϕ σ 2.43±0.08 4.88±0.10 7.29±0.01 9.79±0.15 14.52±0.02 19.39±0.15

1.42 1.46 0.07 1.17 0.13 1.16

On Heating ϕ σ 2.40±0.15 4.79±0.17 7.25±0.03 9.57±0.19 14.42±0.24 19.25±0.38

2.01 2.05 0.25 1.84 2.19 2.32

Part of this variation may be attributed to ‘numerical noise’ introduced during the commutation of φ. Figure 3 shows typical plots. Average real heating and cooling rates were found to be 2 to 4% lower than the nominal values. Considering that typical uncertainties in DSC measurements are of the order of 5%, the discrepancies found between nominal and ‘true’ heating/cooling rates are not a serious problem[19].

3. Results and Discussion 3.1. DSC measurements The DSC scans for heating, cooling and reheating stages, obtained according to the experimental procedure described above, are shown in Figure 4. Figure 4a presents the endotherms corresponding to the first melting of PHB, which takes place as double melting peaks, and may be affected by the injection molding process to which specimens were submitted. Figure 4b shows a crystallization event during cooling from the melt, in the form of broad peaks, particularly at low cooling rates. In addition to the second melting endotherms, Figure 4c shows a previous cold crystallization event (the exothermic peak) at all but the slowest heating/cooling rate (2.5 °C/min).

Thus, under experimental conditions tested the first melting occurs as double melting peaks (a major peak, followed of a minor peak at higher temperature) and the second fusion is visualized as a single complex peak (with a shoulder – hidden peak – at a lower temperature). The crystallization occurs partially during the cooling stage (from the melt) and partially during the reheating stage (as cold crystallization).

The relative areas of melt and cold crystallization peaks (related to the amount crystallized) depend on the rate during the cooling stage. No cold crystallization was detected at rates below 5°C/min (see Table 3); it is probable that during cooling stage at φ < 5 °C/min the all crystallizable polymer effectively crystallized from the melt.

Figure 3. Typical real heating and cooling rates; cooling at nominal 5°C/min (a), heating at nominal 15°C/min (b). Horizontal black lines represent the average (real) heating/cooling rate in each case.

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Melting and crystallization of poly(3-hydroxybutyrate): effect of heating/cooling rates on phase transformation 3.1.1. Crystallization analysis Fractional crystallization versus time plots are presented in Figure 5 for the cooling/reheating rates tested, and some thermodynamic and kinetic parameters of both melt and cold crystallization are presented in Table 2.

The crystallization curves show the sigmoid shape characteristic of phase transformation processes in polymers.

Macroscopic crystallization in polymers starts at very low rates due to the first, slow nucleation step; crystallization rate increases during the main or bulk crystallization and then decreases as the material is depleted of crystallizable molecules

and due to spherulitic impingement[20-22]. Most crystallization parameters are strongly dependent on the cooling/reheating rate, and the dependence is different for the melt and cold crystallization processes. The crystallization temperature increases with the heating rate for the cold crystallization, while the opposite trend is observed for crystallization from the molten state. Moreover, cold crystallization is significantly faster than melt crystallization (notice the different time scale in Figures 5b and 5c).

Results presented in Figure 5 and Table 3 evidence differences in phase transformation processes for cold

Figure 4. DSC scans obtained during the heating (a), cooling (b) and reheating (c) stages.

Table 3. Crystallization parameters determined during cooling and reheating stages. ϕNOM

ϕ

(°C/min) 2.5 2.4 5 4.9 7.5 7.3 10 9.7 15 14.5 20 19.5

TC

(°C) 83.9 62.0 64.5 56.6 60.0 56.4

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Melt crystallization (C1) ΔTC ΔHC cMAX

(°C) 37.0 54.1 53.6 31.8 51.7 61.0

(J/g) 49.8 40.3 39.9 13.8 5.4 5.5

(min−1) 0.114 0.528 0.228 0.528 0.414 0.504

τ½

(min) 9.28 7.50 4.70 1.60 2.20 1.93

TC

(°C) – 52.2 55.8 58.5 67.1 73.3

Cold crystallization (C2) ΔTC ΔHC cMAX

(°C) – 34.0 19.6 24.9 30.9 34.1

(J/g) – 5.7 6.7 12.5 32.8 34.0

(min−1) – 0.510 0.642 0.732 0.990 1.170

τ½

(min) – 2.68 1.95 1.77 1.67 1.38

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Wellen, R. M. R., Rabello, M. S., Araujo, I. C., Jr., Fechine, G. J. M., & Canedo, E. L. crystallization (amorphous solid → crystalline solid) and melt crystallization (fluid → crystalline solid). Figure 6 shows graphically some of these trends.

Table 4 compares the crystallinity change determined during the cooling ∆Xc (C1) and reheating ∆Xc (C2) stages, as well as the total crystallinity change ∆Xc = ∆Xc (C1) + ∆Xc (C2). The fraction of crystallizable polymer that crystallizes in each stage is a strong function of the cooling/heating rate.

Figure 7 shows these results in graphical form.

The macroscopic trends revealed are consistent with moderately fast crystallization kinetics, overrun on cooling: it appears that at the faster cooling rates the material doesn’t have enough time to complete crystallization before molecular motions slow down at low temperatures.

Considerations of the microstructure implication of these trends require a microkinetic study[23-26] beyond the scope of this paper. However, the results presented in Tables 3, 4 and Table 4. Crystallization parameters determined during cooling and reheating processes. ϕNOM 2.5 5.0 7.5 10.0 15.0 20.0

ΔXC (C1) 33.5 27.6 27.3 9.5 3.7 3.8

ΔXC (C2)

ΔXC

(oC/min) (%) <0.5 33.5 3.9 31.5 4.6 31.9 8.6 18.0 22.5 26.2 23.3 27.1

ΔXC(C2)/ΔXC 0 0.124 0.144 0.475 0.859 0.861

Figure 5. Fractional crystallization x (%) versus crystallization time t (min) for the melt crystallization (C1) during the cooling stage (a) and for the cold crystallization (C2) during the reheating stage (b).

Figure 6. Crystallization temperature (a) and maximum crystallization rate (b) in terms of the cooling/heating rate, for the melt and cold crystallization processes.

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Melting and crystallization of poly(3-hydroxybutyrate): effect of heating/cooling rates on phase transformation Figures 6, 7 suggest that controlling the cooling/heating rate may result in crystals grown at different temperature intervals, perhaps with different stabilities and microstructures. From the literature it is known that the microstructure (the size and perfection of the crystalline entities), as well as the ratio of crystalline/amorphous phases, have direct influence on the properties of polymeric materials[20,27-29].

3.1.2. Melting analysis A complex peak structure was observed for the second melting event: a major peak preceded (and sometimes succeeded) by minor peaks that may be reduced to “shoulders” (hidden peaks) on the main peak. Multiple melting peaks may be attributed to the melting of different types of crystals, with different sizes and thermal stabilities. Specifically, small and less perfect crystals melt at lower temperature

Figure 7. Fraction of crystallizable PHB that cold crystallizes, as a function of cooling/heating rate.

than the larger and more perfect ones. Multiple melting peaks observed by DSC are a rather common feature for many semi crystalline polymers, including polyesters.

Multiple peaks may be attributed also to the existence of different crystal modifications, or to melting-recrystallization processes occurring during the DSC scan[26,30-33].

From DSC scans of Figure 4, the molten fraction xF for the second fusion (F2) during the reheating stage was computed by integration of the endothermic peaks.

Figure 8 shows the plots of molten fraction as a function of time and temperature. Melting curves are also sigmoid; however, imperfect probably due to deformations caused the secondary peaks and shoulders observed in Figure 4c.

Table 5 presents some parameters determined during second melting event of the PHB, including the melting peak temperature Tmp and the temperature for 99% completion of melting Tm (which may be considered the true observed melting point of the resin, according to literature[34] the temperature interval for the melting process ∆Tm, the latent heat of melting per unit mass ∆Hm, and the crystallinity (∆Xc). Two measures of the kinetics of melting are included: the maximum melting rate cmax (the melting rate at the peak temperature) and the half-melting time τ½ (the time required to melt one-half of the polymer that melts, which is inversely proportional to average melting rate between xF = 0 and xF = 0.5).

A moderate increase of the melting temperature with the heating rate was observed, especially for the higher heating rates, and – consistently with the previous observations – a significant, almost linear, increase of the rate of fusion with increasing heating rate. Figure 9 shows these results in graphical form. The data in Tables 4 and 5 show that the total amount of crystallinity developed during the cooling/reheating cycles decreases at high cooling/heating rates (φnom > 7.5°C/min).

Figure 8. Evolution of molten fraction with time (a) and temperature (b) for the melting event during the reheating stage (F2).

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Wellen, R. M. R., Rabello, M. S., Araujo, I. C., Jr., Fechine, G. J. M., & Canedo, E. L. Table 5. Melting parameters determined during reheating process. ϕNOM

ϕ

(°C/) 2.5 5 7.5 10 15 20

TMP

TM

ΔTM

(°C) 2.4 4.9 7.3 9.7 14.5 19.5

168 168 168 169 171 173

172.4 171.8 183.5 178.5 179.0 182.0

23.0 36.4 47.4 42.1 58.8 35.5

ΔHM

XC

cMAX

τ½

(J/g)

(%)

(min−1)

(min)

59.9 66.6 88.0 57.2 62.4 50.5

41.0 45.6 60.3 39.2 42.7 34.6

0.294 0.474 0.522 0.822 1.002 1.398

8.90 7.73 5.40 4.68 4.40 1.70

Figure 9. Some measures of the thermodynamics and kinetics of melting determined during the reheating stage.

4. Conclusions

6. References

This paper presents a methodological procedure to study nonisothermal crystallization and melting phenomena in PHB by differential scanning calorimetry. The crystallization and melting behavior of PHB is strongly affected by the rates of cooling and heating. Increasing the cooling rate decreases the melt crystallization temperature and the amount of polymer that crystallizes from the molten state, and increases the rate of crystallization and the amount of polymer that cold crystallizes upon reheating. Increasing the reheating rate increases the cold crystallization temperature and the rate of crystallization from the amorphous solid phase, the total crystallinity developed is also affected, decreasing at high rates of cooling/heating within the interval studied.

These findings suggest that the crystallinity of PHB may be controlled by the thermal cycles of heating and cooling to which the material is subjected.

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15. Canetti, M., Urso, M., & Sadocco, P. (1999). Influence of the morphology and of the supermolecular structure on the enzymatic degradation of bacterial poly(3-hydroxybutyrate).

Polymer, 40(10), 2587-2594. http://dx.doi.org/10.1016/S00323861(98)00503-5.

16. El-Hadi, A., Schnabel, R., Straube, E., Müller, G., & Riemschneider, M. (2002). Effect of melt processing on crystallization behavior and rheology of poly(3-hydroxybutyrate) (PHB) and its blends.

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18. Barham, P. J., Keller, A., Otun, E. L., & Holmes, P. A. (1984).

Crystallization and morphology of a bacterial thermoplastic: poly-3-hydroxybutyrate. Journal of Materials Science, 19(9), 2781-2794. http://dx.doi.org/10.1007/BF01026954.

19. Höhne, G. W. H., Hemminger, W. F., & Flammersheim, H. J. (2010). Differential scanning calorimetry (2nd ed.). Berlin: Springer.

20. Wellen, R. M. R., & Rabello, M. S. (2005). The kinetics of isothermal crystallization and tensile properties of poly(ethylene) terephthalate. Journal of Materials Science, 40(23), 6099-6104. http://dx.doi.org/10.1007/s10853-005-3173-3.

21. Ziaee, Z., & Supaphol, P. (2006). Non-isothermal melt- and cold-crystallization kinetics of poly(3- hydroxybutyrate).

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22. Wellen, R. M. R., & Rabello, M. S. (2009). Antinucleating action of polyestyrene on the isothrermal cold crystallization

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23. Shafee, E. E., & Ueda, W. (2002). Crystallization and melting behavior of poly(ethylene oxide)/poly(n-butyl methacrylate) blends. European Polymer Journal, 38(7), 1327-1335. http:// dx.doi.org/10.1016/S0014-3057(02)00018-6.

24. An, Y., Dong, L., Mo, Z., Liu, T., & Feng, Z. (1998). Nonisothermal crystallization kinetics of poly(β-hydroxybutyrate). Journal of Polymer Science: Part B, Polymer Physics, 36(8), 1305-1312. http:// dx.doi.org/10.1002/(SICI)1099-0488(199806)36:8<1305::AIDPOLB5>3.0.CO;2-Q. 25. Wellen, R. M. R., & Rabello, M. S. (2007). Redução da velocidade de cristalização a frio do PET na presença de poliestireno. Polímeros: Ciência e Tecnologia, 17(2), 113-122. http://dx.doi.org/10.1590/S0104-14282007000200010.

26. Wellen, R. M. R., Canedo, E. L., & Rabello, M. S. (2011).

Nonisothermal cold crystallization of poly(ethylene terephthalate).

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27. Blundell, D. J., Liggat, J. J., & Flory, A. (1992). The crystal lamellar morphology of an aromatic polyketone with unusual crystallization and melting behavior. Polymer, 33(12), 24752482. http://dx.doi.org/10.1016/0032-3861(92)91127-N. 28. Blundell, D. J., & Osborn, B. N. (1983). The morphology of poly(aryl-ether-ether-ketone). Polymer, 24(8), 953-958. http:// dx.doi.org/10.1016/0032-3861(83)90144-1.

29. Varma, P., Lofgren, E. A., & Jabarin, S. A. (1998). Properties and kinetics of thermally crystallized oriented poly(ethylene terephthalate) (PET). II. Physical and optical properties.

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30. Gunaratne, L. M. W. K., & Shanks, R. A. (2005). Melting and thermal history of poly(hydroxybutyrate-co- hydroxyvalerate) using step-scan DSC. Thermochimica Acta, 430(1-2), 183-190. http://dx.doi.org/10.1016/j.tca.2005.01.060.

31. Harding, K. G., Dennis, J. S., von Blottnitz, H., & Harrison, S. T. (2007). Environmental analysis of plastic production processes: comparing petroleum-based polypropylene and polyethylene with biologically-based poly-beta-hydroxybutyric acid using life cycle analysis. Journal of Biotechnology, 130(1), 57-66. http://dx.doi.org/10.1016/j.jbiotec.2007.02.012.

PMid:17400318.

32. Papageorgiou, G. Z., & Karayannidis, G. (1999). Multiple melting behaviour of poly(ethylene-co-butylene naphthalate2,6-dicarboxylate)s. Polymer, 40(19), 5325-5332. http://dx.doi. org/10.1016/S0032-3861(98)00746-0.

33. Renstad, R., Karlsson, S., Albertsson, A. C., Werner, P. E., & Westdahl, M. (1997). Influence of processing parameters on the mass crystallinity of poly(3-hydroxybutyrate- co-3hydroxyvalerate). Polymer International, 43(3), 201-209. http:// dx.doi.org/10.1002/(SICI)1097-0126(199707)43:3<201::AIDPI761>3.0.CO;2-7.

34. Menczel, J. D., & Prime, R. B. (2009). Thermal analysis of polymers. New York: Wiley.

Received: Oct. 11, 2014 Revised: Nov. 29, 2014 Accepted: Jan. 09, 2015

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List of symbols cmax Maximum crystallization/melting rate (at Tc or Tmp) E0 Total latent heat released or absorbed by the sample during the crystallization or melting event J Heat flow; rate of thermal energy exchanged between sample and the surroundings J0 Virtual base line during a phase change event mS Mass of the sample Tc Peak crystallization temperature Tmp Peak melting temperature Tm Melting point (temperature at 99% of the amount molten) τ½ Crystallization/melting half-time (time to attain 50% fractional crystallization/melting), inversely proportional to the mean crystallization/melting rate between 0 and 50% fractional crystallization/fusion x fractional crystallization during a crystallization event xF molten fraction during a melting event φnom Nominal (set) heating/cooling rate φ True mean heating/cooling rate(computed by Equation 7) ∆Tc Crystallization interval (from 1% to 99% of the amount crystallized) ∆Tm Melting interval (from 1% to 99% of the amount melted) ∆Hc Specific latent heat of crystallization ∆Hm Specific latent heat of melting ∆Xc Mass crystallinity (total) ∆Xc(C1) Crystallinity developed during the crystallization form the melt event (C1) ∆Xc(C2) Crystallinity developed during the cold crystallization event(C2)

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http://dx.doi.org/10.1590/0104-1428.1869

Development of paints with infrared radiation reflective properties Eliane Coser1, Vicente Froes Moritz1, Arno Krenzinger2 and Carlos Arthur Ferreira1 Laboratório de Materiais Poliméricos – LAPOL, Programa de Pós-Graduação em Engenharia de Minas, Metalúrgica e de Materiais – PPGE3M, Universidade Federal do Rio Grande do Sul – UFRGS, CEP 91501-970, Porto Alegre, RS, Brazil 2 Laboratório de Energia Solar – LABSOLAR, Programa de Pós-Graduação em Engenharia Mecânica – PROMEC, Universidade Federal do Rio Grande do Sul – UFRGS, CEP 91501-970, Porto Alegre, RS, Brazil 1

*cosereliane@yahoo.com.br

Abstract Large buildings situated in hot regions of the Globe need to be agreeable to their residents. Air conditioning is extensively used to make these buildings comfortable, with consequent energy consumption. Absorption of solar visible and infrared radiations are responsible for heating objects on the surface of the Earth, including houses and buildings. To avoid excessive energy consumption, it is possible to use coatings formulated with special pigments that are able to reflect the radiation in the near- infrared, NIR, spectrum. To evaluate this phenomenon an experimental study about the reflectivity of paints containing infrared-reflective pigments has been made. By irradiating with an IR source and by measuring the surface temperatures of the samples we evaluated: color according to ASTM D 2244-14, UV/VIS/NIR reflectance according to ASTM E 903-12 and thermal performance. Additionally, the spectral reflectance and the IR emittance were measured and the solar reflectance of the samples were calculated. The results showed that plates coated with paints containing IR-reflecting pigments displayed lower air temperature on the opposite side as compared to conventional coatings, indicating that they can be effective to reflect NIR and decrease the temperature of buildings when used in roofs and walls.

Keywords: cool paints, near-infrared reflectance, solar spectral reflectance, cool pigments, colored reflecting pigments.

1. Introduction Solar energy plays an important role in economic development around the world. The wavelength of the light that reaches the Earth’s surface conventionally ranges from 300 to 2500 nm. The human eye is sensitive to only a part of the electromagnetic spectrum. Pigments are colour compounds that are responsible to give colour to objects as they selectively absorb the visible light and reflect the remainder corresponding to its colour. Apart from the visible region, pigments also interact with other wavelengths of light in the electromagnetic spectrum[1-3].

Roughly 5% of the Sunlight that reaches the Earth’s surface is in the form of ultraviolet (UV) (wavelength between 300 and 400nm) (see Figure 1) which is the main responsible by the photo degradation of organic materials including organic coatings. Around 42% of the solar energy occurs in the visible region of the electromagnetic spectrum.

Different colours are detected by the optical human system in the wavelength range from 400 to 700 nm. Some 53% of the total solar energy is in the infrared region (IR) whose wavelength ranges from 700 to 2500nm. Heat is a direct consequence of either visible or infrared radiation incident on an object. The heat-producing region of the infrared radiation ranges from 700 to 1100 nm[1,4].

A roof with high solar reflectance (the ability to reflect sunlight) and high thermal emittance (the ability to radiate heat) remains cool in the sun, reducing demand for cooling power in air-conditioned buildings and increasing occupant

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comfort in unconditioned buildings. Increasing the solar reflectance lowers a surface’s temperature, since solar radiation is reflected rather than absorbed. In turn, this decreases the heat penetrating into the building especially during summer, resulting in more comfortable thermal conditions if the building is not air-conditioned[5,6].

The primary purpose of IR-reflective coatings is to keep objects cooler than they would be using conventional pigments. Normally the reflection of solar energy is empirically obtained by the use of white paints, which produce a beautiful effect but is little useful and acceptable in big cities. Normal paints containing titanium dioxide (TiO2) as white pigments reflect visible and IR radiation very well[7].

TiO2 is the most employed pigment in the formulation of paints; it is an ingredient which improves the quality of paint ensuring high coverage power, durability, brightness and opacity[8]. To avoid the monotonic effect that a totally white city could produce, there are coloured paints on the marked, but the normal pigments used in these paints have low reflectivity. Each pigment has distinct IR-reflective characteristics. In addition to their reflective properties, pigments can differ in their weatherability, chemical resistance, and other durability criteria.

Inorganic pigments infrared reflectors are colour pigments made of inorganic complexes, which reflect the wavelengths in the infrared region as well as selectively reflect visible light. The reflectivity and absorptivity are dependent of the

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Coser, E., Moritz, V. F., Krenzinger, A., & Ferreira, C. A. pigment. Therefore, an infrared-reflective pigment may be in any colour and it can be synthesized by the calcination of a mixture of oxides, nitrates, acetates and even metal oxide at temperatures above 1000°C. At the calcination temperature the solids themselves become reactive, metal and oxygen ions rearrange to form new crystalline stable structure such as rutile or spinel[1,9]. The pigment particle size is of extreme importance to the NIR reflectance. Pigments consisting of smaller particles or nanoparticles significantly improve the reflective properties[10].

Libbra et al.[4] prepared an acrylic water based paint containing a mixture of a red Fe2O3 pigment and TiO2 and applied it on terra-cotta tiles in Italy. They verified that even using a highly IR reflective pigment (TiO2), the reflectance of the protective layer was not satisfactory.

Ryan[11] has evaluated the total solar reflectance (TSR) of commercial cool pigments and Artic pigments from Shepherd Colour Company[12] with different colours.

They verified that mixing the Blue 211 with the Black 10C909, the reflectance decreased when he increased the amount of the black pigment. He recommended that formulators must be extremely careful during mixing the components of a paint in order to maintain IR reflectance, as even small amounts of IR-absorbing pigments can strongly reduce TSR. An extensive study of cool pigments has been conducted by Malshe and Bendiganavale[1]. They tried to produce IR reflective pigments similar to the commercials by calcination of mixtures of several oxides of metals as Ti, Co, Ni and Mn in different atomic proportions. None of their experiments resulted in significant IR reflectivity, demonstrating the complexity of the subject. They have used a very simple device to evaluate the difference of temperature attained by painted samples irradiated by an IR Lamp. Uemoto et al.[13] have used three commercial paints formulated by an industrial collaborator. The paints were similar in colour and have been applied on cement roof sheets. They measured the difference of temperature using a closed wood device and the results demonstrated a higher NIR reflectance for samples coated with cool paints compared to those coated with conventional paints.

New coatings that are coloured but still reflect sunlight and remain cooler are being developed using specialized cool pigments. These smart cool coatings decrease roof surface temperature, reducing the energy needed for cooling buildings and making unconditioned buildings more comfortable.

In the present work a procedure similar to[1] has been adopted. In order to demonstrate that colours coatings can reflect sunlight and heat in a similar way to white coatings, paints in four colour were formulated at the laboratory with cool and conventional pigments and applied to fibre cement plates. The samples were characterised by measuring the colour, the spectral reflectance and the temperature in front and behind the samples. The diffuse reflectance and infrared emittance were measured and the solar reflectance of the samples was calculated.

2. Experimental 2.1. Preparation of paints The materials were classified according to their thermal performance and physical properties in cool and warm materials. Eight samples were analysed: four cool paints tested in comparison to four samples similarly coloured prepared using conventional pigmented. The solvent-based paints (bi-component) were prepared in the laboratory using a Byk model Dispermat N1 disperser with a Cowles disk according to the basic formulation shown in Table 1, changing only the pigment. The materials were applied with a brush, in two layers. The final dry thicknesses were 75 +/–µm. Each paint was tested in at least three samples.

The pigments used in the production of reflective paints were provided by Shepherd Colour Company representative office in Brazil. In Table 2 are described the pigments and the chromospheres that each belongs pigment used for the production of paint[12].

All the coatings were applied on 20 x 20 x 0.8 cm fibre cement plates. The aspect of the samples is shown in Figure 2.

Table 1. Formulation of paints produced in laboratory.

COMPONENT 1

PARTS (%)

Resins

Hydroxylated acrylic

50.00

Solvent

Organic: Ethyl glycol acetate, Ethyl acetate and Xylene

27.50

0.20

Thickener

Tixogel

Dispersant

Additive BYK 108

0.30

Pigment

TiO2

22.00

Total

100.00

COMPONENT 2

PARTS (%)

Catalyst

Isocyanate AQ - 6008

Solvent

Organic Total

38.00 62.00 100.00

Table 2. Description of the pigment.

Figure 1. The solar spectrum[5].

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Pigment

Chromophores

Yellow 346

Chrome Antimony

Description samples Yellow

Brown 157

Zinc Iron Chromite

Brown 1

Brown Rosse 208

Iron oxide

Brown 2

Black 28

Copper Chromite

Black

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Development of paints with infrared radiation reflective properties 2.2. Characterization of coatings A UV/Vis/NIR spectrophotometer (Varian Carry 5000) was used for measure the spectral reflectance of the samples[14]. The spectrophotometer was fitted with a 150 mm diameter integrating sphere (Labsphere DRA 1800) which collects both specular and diffuse radiation. The reference reflectance material used for the measurement was a PTFE plate, according to ASTM Standard E 903-12[5,15].

Colours of samples were measured according to the CIE (Commission Internationale de l’Eclairage) that is the regulatory body responsible for international recommendations for photometry and colorimetry[16], which is being presently widely used to characterize the colour of the paints that were produced in the laboratory (reflective paint) compared with the commercial paints.

In this system, the values measured are L*, a* and b* and are called CIELAB. Each colour is represented by the coordinates in a three-dimensional system generating a set of three members (Figure 3).

The vertical L* axis represents the differences between light (L*=100) at the top and dark (L*=0) at the bottom.

The axis a* displays the difference between red (+a*) and green (–a*), while b* corresponds to the difference between yellow (+b*) and blue (–b*).

Colour difference (ΔE*) between two colour points in the CIELAB space are calculated as the Euclidean distance between their locations in the three-dimensional space defined by L*, a*, and b*[18] . Thus, mathematically, it is calculated using the formula[16]: ∆E =

( ∆L* ) + ( ∆a* ) + ( ∆b* ) 2

2

2

Two thermocouples were used to record the temperatures, one thermocouple in front and the other at the back of the plate. The plates were allowed under infrared lamp during an hour to equilibrate the temperature. After this period the temperature was recorded. The same procedure was adopted for the uncoated sample. The difference between the temperature in front (coated) and in the back (uncoated) side of the samples was used as an indicative of the performance of the pigments (infrared-reflective or not). Figure 4 shows the experimental setup for simulating the infrared reflectivity of pigments.

The temperature of the room was kept constant in order to avoid accumulation of heat in this space and to assure that the heat absorbed by the fibre cement plates was exclusively by irradiation and not by convection of the air in contact with the inner surface of the plates.

The solar reflectance (ῤ) and the thermal emittance (Ɛ)[20] of a roof surface are important surface properties affecting the roof temperature, which, in turn, drives the heat flow through the roof[21,22]. The emissivity of the samples was also measured using the device model AE Emissometer Devices & Services, according to ASTM standard C1371-10.

(1)

Measurements of colour were done in dry films according to ASTM D 2244-14[19], using a Byk Gardner model Spectro‑guide Portable Spectrophotometer.

The painted plates were exposed to a Philips PRA 38 IR Red (150W 230V reference E27 ES) infrared lamp. The distance between the sample and the lamp was maintained at 30 cm.

Figure 3. CIE colour system L*, a*, b*[17].

Figure 2. Samples painted with cool and conventional pigments with visual similar colour.

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Figure 4. Experimental setup for measuring temperature as IR lamps irradiates coated cement plates.

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3. Results and Discussion 3.1. Colour characterization of paint films Although visually the samples of the same colour were quite similar (Figure 2), the CIELAB results presented in Table 3 display significant differences between cool colour paints and standard paints. These values indicate that the lower L* of the paint, the lower is its luminosity[18,23].

L* values indicate the difference in luminosity between samples. The data obtained show that standard sample presented lower luminosity than cool sample.

According to Table 3, L* of Yellow paint is higher for the reflective sample than for standard paint. Despite visually having the same colour, as the b* values are positive

and high, ΔE indicates the existence of a colour difference between these Yellow coloured paints.

Standard and cool Black paints are very similar, as ΔE is 0.85 ± 0.10 (it would be zero if the colours were exactly the same). The Brown 1 paints presented ΔE=1.51 ± 0.18, but displayed similar lightness L* values (37.8 and 38.2).

The Brown 2 paint displayed quite different colours (ΔE=2.51 ± 0.30) and this paint is more reddish than Brown 1 because a* parameter of these sample (a*~ 18-19) are higher than brown 2 sample (a*~ 27-29).

3.2. The UV, VIS and NIR reflectance The spectrophotometric reflectance spectra (UV, VIS, and NIR) of all samples are shown in Figure 5. Measurements were done according to ASTM E 903-12. The reflectance of

Table 3. Colour characterization of standard and cool paints.

Standard

L* 61.74±0,32

Colour coordinates a* 20.62±0,18

b* 52.52±0,22

Yellow

64.35±0,16

22.64±0,15

53.19±0,15

Cool Yellow Standard Black

28.48±0,17

–0.75±0,05

–0.99±0,10

0.85 ± 0.10

Cool Black Standard Brown 1

28.11±0,22 37.84±0,11

-0.85±0,07 18.44±0,11

–0.72±0,08 16.44±0,05

1.51 ± 0.18

Cool Brown 1 Standard Brown 2

38.23±0,04 41.20±0,28

19.65±0,12 27.61±0,09

17.73±0,11 22.84±0,20

2.51 ± 0.30

Cool Brown 2

44.63 ±0,06

29.75±0,16

23.15±0,015

Sample

∆E 3.05 ± 0.37

Figure 5. Reflectance of paints in the UV/Vis/NIR region (a) Yellow, (b) Brown 1, (c) Brown 2 and (d) Black.

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Development of paints with infrared radiation reflective properties Table 4. Temperature in front and behind painted plates.

Sample Uncoated Standard

Front of Panel (°C) 51,3 ± 0.21 56.2 ± 0.77

Behind Panel (°C) 40,8 ± 0,12 40.3 ± 0.78

∆T1 (°C) 10.5 15.7

∆T2 (°C)

Yellow Cool

51.7 ± 0.3

39.4 ± 0.10

12.3

–0.9

Yellow Standard Black Cool Black Standard Brown 1 Cool Brown 1 Standard Brown 2 Cool Brown 2

61.6 ± 0.07 59.5 ± 0.14 60.3 ± 0.70 54.3 ± 1.20 58.6 ± 0,07 53.9 ± 0.08

45.8 ±0.20 45.3 ± 0.15 40.3 ± 0.21 41.3 ± 0.85 44.6 ± 0,07 41.9 ± 0.07

15.8 14.3 20.0 13.1 14.0 12.0

–0.5 +1.0 –2.7

∆T1: Tfront - Tbehind; ∆T2: Tbehind cool - Tbehind standard.

all points sample is very low at the UV region (around 5%).

As expected, Figures 5a, b and c show that the cool paints presented higher reflectance than standard paints in the VIS region, namely between 2-3% at 400nm and 10‑40% for standard and 28-60% for cool paints at 700nm. Black paint (Figure 5d) present very low VIS reflectance.

The NIR spectra display remarkable differences. Yellow and Brown 1 reflectance spectra of paints prepared with cool and conventional pigments (Figure 5a and Figure 5b) show differences that reach some 35% (Yellow) and 40% (Brown1) demonstrating clearly the reflective proprieties of cool pigments. Brown 2 cool paint (Figure 5c) is less reflective than Yellow and Brown1 cool paint. The Brown 2 cool and standard paints showed reflectance that are closer than for the other samples. These results show clearly that Brown 2 cool pigment is less efficient in reflect VIS and NIR radiation than Yellow and Brown 1 pigment.

This high NIR “invisible” reflectance displayed by the samples explains the fact that the cool colour are characterized by high reflectance values. Sunlight is more intense in the visible region, but it also emits a substantial amount of energy in the invisible ultraviolet (UV) and near-infrared (NIR).

In fact, about half percent of all solar power that reaches the Earth’s surface arrives as invisible near-infrared radiation.

The reflective ability of the samples to reflect IR radiation was tested by measuring the temperature in front of and behind painted fibre concrete plates using the device shown in Figure 3. The Table 4 shows the results of the temperatures measured for the coatings.

It is clear from Table 4 that even uncoated plates made with fibre cement are able to display different temperatures on the two sides if they are irradiated by an IR lamp.

The temperatures behind the panels for all cool pigments are lower than the temperatures for conventional pigments, except for Brown 1 (see also ∆T1). This behaviour is the consequence of the ability of cool pigmented coatings to reflect IR radiation from the lamp, reducing the temperature behind the panels. For Brown 1 ∆T1 is the highest among all paints, even if the temperature in front of the panel is the highest of all sample.

As a consequence, the lowest temperature behind the panels has been measured to the yellow cool sample (39.4 °C) and the higher temperature to the standard Black sample (45.8°). Yellow pigment demonstrated to be a good Polímeros, 25(3), 305-310, 2015

Table 5. Solar reflectance and emissivity of paints.

Sample

Emissivity

Cool Yellow Standard Yellow Cool Black Standard Black Cool Brown 1 Standard Brown 1 Cool Brown 2 Standard Brown 2

0.92 0.87 0.90 0.91 0.90 0.92 0.89 0.89

Solar reflectance 0.52 0.31 0.06 0.05 0.29 0.85 0.24 0.18

choice to reflect IR radiation in external roof of buildings.

As another form to express the behaviour of cool paints, the difference of temperature behind the panels coated with cool and standard paints (∆T2) have been calculated and the cool Brown 2 sample showed the higher value (–2.7 °C) indicating it is also a good choice to external roof colour that reflect IR radiation.

The emissivity is the measure of an object’s ability to emit infrared energy. Emitted energy indicates the temperature of the object. Emissivity can have a value from 0 (shiny mirror) to 1.0 (Blackbody). The emissivity (ε) is a surface characteristic of each material. The higher the emissivity, higher the heat loss by the substrate. The emissivity (ε) of all paints is in the range 0.87 to 0.92 (Table 5), indicating that all paints have an appropriate behaviour to be used as a roof coating as they are able to loose heat.

The solar reflectance results shown in Table 5 were calculated according to the standard ASTM G173-12.

They were obtained by the integral of the data in Figure 5 with reference to Solar Spectral Irradiances.

When applying an infrared reflective coating or cool coating, one must maximize solar reflectance and emissivity, and minimize all contamination by infrared absorbing materials. It is observed in Table 5 that the commercial coatings have lower reflectance than cool coatings, with the exception of the coating Standard Brown 1.

The results displayed at Table 5 show that coatings prepared with cool pigments emit infrared radiation more efficiently than the standard coatings.

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4. Conclusions Eight types of coatings were studied, four coatings containing cool pigments and four coatings containing conventional pigments. Their colours were very similar, as measured by the CIE standard. The use of cool pigments in paint formulations allows the development of coatings with similar reflection in the colours (VIS range) but higher reflectance in NIR radiation. Cement fibre plates coated with cool paints displayed lower temperature than standard coated panels when irradiated by an IR lamp. Consequently, if roofs and walls are coated with paints containing cool pigments, the temperature inside the buildings and houses could be maintained lower than if they were painted with conventional paints. In this way, the cool-coloured paints can be employed as an alternative to white paint and can improve thermal comfort conditions of low-cost housing, industrial and residential buildings constructed with fibre‑cement roofing sheets.

5. Acknowledgements This work was supported by CNPq. The authors would like to thank Águia Química S.A. for providing the resin and Color Net for providing the pigments used. Joe Tonolli is thanked for preparing the matched standard coloured paints used in this work.

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2. Levinson, R., Berdahl, P., & Akbari, H. (2005). Solar spectral optical properties of pigments— art I: model for deriving scattering and absorption coefficients from transmittance and reflectance measurements. Solar Energy Materials and Solar Cells, 89(4), 319-349. http://dx.doi.org/10.1016/j.solmat.2004.11.012.

3. Kaur, B., Quazi, N., Ivanov, I., & Bhattacharya, S. N. (2012).

Near-infrared reflective properties of perylene derivatives. Dyes and Pigments, 92(3), 1108-1113. http://dx.doi.org/10.1016/j. dyepig.2011.06.011.

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Spectral response data for development of cool coloured tile coverings. Optics & Laser Technology, 43(2), 394-400. http:// dx.doi.org/10.1016/j.optlastec.2009.07.001.

5. Levinson, R., Akbari, H., & Berdahl, P. (2010). Measuring solar reflectance—Part I: Defining a metric that accurately predicts solar heat gain. Solar Energy, 84(9), 1717-1744. http://dx.doi. org/10.1016/j.solener.2010.04.018.

6. Zinzi, M., Carnielo, E., & Agnoli, S. (2012). Characterization and assessment of cool coloured solar protection devices for Mediterranean residential buildings application. Energy and Building, 50, 111-119. http://dx.doi.org/10.1016/j.enbuild.2012.03.031.

7. Song, Z., Zhang, W., Shi, Y., Song, J., Qu, J., Qin, J., Zhang, T., Li, Y., Zhang, H., & Zhang, R. (2013). Optical properties across the solar spectrum and indoor thermal performance of cool white coatings for building energy efficiency. Energy and Building, 63, 49-58. http://dx.doi.org/10.1016/j.enbuild.2013.03.051.

8. DuPont (2007). DuPont™ Ti-Pure titanium dioxide: titanium dioxide for coatings. U.S.A.: DuPont. Retrieved in 15 May 2014, from http://www2.dupont.com/Titanium_Technologies/pt_US/ tech_info/literature/Coatings/CO_B_H_65969_Coatings_Brochure. pdf

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11. Ryan, M. (2005). Introduction to IR-reflective pigments. Paint & Coatings Industry, 170-176.

12. Shepherd Color Company (2001). Select a product for more information. Cincinnati: Shepherd Color Company. Retrieved in 15 May 2014, from http://www.shepherdcolor.com/Products/ ColorChart.aspx 13. Uemoto, K. L., Sato, N. M. N., & John, V. M. (2010). Estimating thermal performance of cool colored paints. Energy and Building, 42(1), 17-22. http://dx.doi.org/10.1016/j.enbuild.2009.07.026.

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West Conshohocken: ASTM. 16. Khan, M. A. I., Ueno, K., Horimoto, S., Komai, F., Someya, T., Inoue, K., Tanaka, K., & Ono, Y. (2009). CIELAB color variables as indicators of compost stability. Waste Management, 29(12), 2969-2975. http://dx.doi.org/10.1016/j.wasman.2009.06.021.

PMid:19781930.

17. Quimanil (2010). Colorimetria. São Paulo: Quimanil. Retrieved in 7 October 2014, from http://www.quimanil.com.br/empresa/ informacoes_detalhe.php?id=7 18. García-Marino, M., Escudero-Gilete, M. L., Heredia, F. J., Escribano-Bailón, M. T., & Rivas-Gonzalo, J. C. (2013). Colorcopigmentation study by tristimulus colorimetry (CIELAB) in red wines obtained from Tempranillo and Graciano varieties.

Food Research International, 51(1), 123-131. http://dx.doi. org/10.1016/j.foodres.2012.11.035.

19. American Society for Testing and Materials (2014). ASTM D2244-14: Standard Practice for Calculation of Color Tolerances and Color Differences from Instrumentally Measured Color Coordinates. West Conshohocken: ASTM. 20. Santamouris, M., Synnefa, A., & Karlessi, T. (2011). Using advanced cool materials in the urban built environment to mitigate heat islands and improve thermal comfort conditions.

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Received: July 29, 2014 Revised: Oct. 21, 2014 Accepted: Jan. 15, 2015

Polímeros , 25(3), 305-310, 2015

http://dx.doi.org/10.1590/0104-1428.1835

Bioactivity, biocompatibility and antimicrobial properties of a chitosan-mineral composite for periodontal tissue regeneration Andrew Paul Hurt1, Arun Kumar Kotha1, Vivek Trivedi1 and Nichola Jayne Coleman1 School of Science, University of Greenwich, ME4 4TB, Chatham Maritime, Kent, United Kingdom

1

*nj_coleman@yahoo.co.uk

Abstract A composite membrane of the polymer, chitosan, and the silver-exchanged mineral phase, tobermorite, was prepared by solvent casting and characterised by scanning electron microscopy and Fourier transform infrared spectroscopy.

The in vitro bioactivity, cytocompatibility and antimicrobial activity of the composite were evaluated with respect to its potential application as a guided tissue regeneration (GTR) membrane. The in vitro bioactivity was verified by the formation of hydroxyapatite on the surface of the membrane in simulated body fluid and its cytocompatibility was established using MG63 human osteosarcoma cells. The presence of silver ions conferred significant antimicrobial activity against S. aureus, P. aeruginosa and E. coli. The findings of this investigation have indicated that the chitosansilver-tobermorite composite is a prospective candidate for GTR applications.

Keywords: chitosan, tobermorite, silver, bioactive, antimicrobial, guided tissue regeneration, periodontal repair.

Introduction Periodontitis - inflammation and progressive destruction of the tooth attachment apparatus - is one of the most widespread infectious diseases in the world[1]. This disease is initiated by the accumulation of bacterial plaque biofilms causing gingivitis, which, if left untreated, leads to the detachment of the epithelial tissue, disconnection of the periodontal ligament (PDL) and the loss of cementum and alveolar bone[1-3]. Traditional treatment of this condition involves the removal of plaque and calculus from the root surface without an attempt to restore the surrounding tissues. The rapidly growing epithelial cells subsequently propagate alongside the tooth root and prevent the re-establishment of the PDL and alveolar bone. Thus the treated periodontal tissues possess a long junctional epithelium and are vulnerable to further bacterial infection and recurrent episodes of periodontitis.

The use of biocompatible membranes for the guided tissue regeneration (GTR) of compromised periodontal structures is an increasingly popular option for the treatment of periodontitis[1-4]. The GTR approach involves the placement of a membrane to exclude the fast-growing soft tissues from the exposed root surface in order to enable the more slow‑growing PDL and hard tissues to regenerate[1-4].

These GTR barriers can be fabricated from either non-resorbable materials (e.g. PTFE) which require surgical removal after service, or resorbable materials (e.g. poly(urethane)) which will biodegrade in situ. In both cases, the principal cause of failure of GTR membranes is bacterial biofilm formation leading to infection; hence, improvements in the design of GTR membranes to incorporate antimicrobial components are now required[1].

Chitosan is the partially N-deacetylated derivative of chitin, a naturally abundant structural polysaccharide obtained from the shells of crustaceans[5-7]. It is non-toxic,

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biodegradable, biocompatible, bioactive, non-antigenic, antimicrobial and finds current biomedical application in wound dressings and medical textiles[5-9]. It is a basic linear co-polymer of glucosamine and N-acetylglucosamine whose structure resembles that of bone extracellular matrix (ECM) [10,11] . In comparison with traditional treatment, the application of bioresorbable chitosan-based GTR membranes has been shown to successfully promote the regeneration of the tooth attachment tissues in canine models[12-14].

Despite its many advantages, chitosan in GTR applications is not sufficiently bioactive to stimulate optimum bone tissue regeneration. It is also prone to bacterial biofilm formation.

Previous studies have indicated that blending chitosan with finely divided bioactive and antibacterial phases, such as bioactive glasses, calcium phosphate minerals and silver nanoparticles can enhance bone tissue regeneration and reduce bacterial biofilm formation[11,15,16].

Tobermorite, Ca5Si6O16(OH)2.4H2O, is a bioactive calcium silicate phase that can be ion-exchanged with antibacterial silver ions[17,18]. The layered structure of tobermorite resembles the calcium silicate hydrate gel phases of the endodontic cements, MTA and Biodentine, which are reported to stimulate the regeneration of cementum and bones tissues[19-21]. A recent in vitro study has demonstrated that the incorporation of synthetic tobermorite into chitosan films enhances both bioactivity and biocompatibility[22].

In this work, a candidate composite GTR membrane of chitosan and silver-exchanged tobermorite has been prepared by solvent-casting and characterised by scanning electron microscopy (SEM) and Fourier transform infrared spectroscopy (FTIR). The bioactivity of the composite membrane was evaluated in vitro and its biocompatibility was assessed using human MG63 osteosarcoma cells. An inhibition zone

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S S S S S S S S S S S S S S S S S S S S

Hurt, A. P., Kotha, A. K., Trivedi, V., & Coleman, N. J. assay was used to determine the antimicrobial activity of the membrane against Staphylococcus aureus, Pseudomonas aeruginosa and Escherichia coli.

Materials and Methods 2.1. Preparation and characterisation Chitosan (50-190 kDa molecular weight, 75-85% deacetylated) and all other reagents were purchased from Sigma-Aldrich, UK, and used without further modification.

Triplicate samples of silver-exchanged tobermorite were prepared characterised by powder X-ray diffraction (XRD) analysis and X-ray fluorescence spectroscopy (XRF), as described in reference[18]. The Bragg peaks of the silverexchanged tobermorite were identified and indexed using JCPDS file 45-1480. The formulae of the original tobermorite and Ag-tobermorite used in this study were Ca4.55Na0.44Si6.00O16.77.6.4H2O and Ca4.06Na0.04Ag0.92Si6.00 O16.54.11.9H2O, respectively.

Chitosan-silver-tobermorite composite membranes (CTAg) were prepared at a chitosan:Ag-tobermorite mass ratio of 10:7 by solvent casting. 2% (w/v) solutions of chitosan were prepared in triplicate by dissolving 0.6 g of the polymer in 30 cm3 of 2% (v/v) acetic acid solution.

0.42 g of Ag-tobermorite were then added and the mixtures were stirred for 5 hours, cast onto polycarbonate plates and dried in air at 60 °C to constant mass to produce the CTAg membranes. Control membranes (labelled CT) for the antimicrobial inhibition zone assay were prepared similarly with the original tobermorite samples.

The CTAg membranes were characterised by FTIR using a Perkin Elmer Paragon spectrometer and by SEM using uncoated samples attached to carbon tabs on an Hitachi SU8030 scanning electron microscope. Secondary electron images were obtained at an accelerating voltage of 1 kV.

CTAg and CT membranes were placed in the centre of each spread plate. Each assay was conducted in quadruplicate.

The plates were examined for clear zones after incubation at 37 °C for 24 hours. The final population densities of the plates spread with S. aureus, E. coli and P. aeruginosa were approximately 1.0 x 109, 5.9 x 108, and 5.9 x 109 colony forming units per plate.

2.4. In vitro biocompatibility of CTAg membrane The in vitro biocompatibility of the CTAg membrane was evaluated using MG63 human osteosarcoma cells (ECACC code: 86051601) as described in reference [22]. In triplicate, either 0, 1 or 4 CTAg membrane strips (1 mm x 4 mm) were incubated with the MG63 cells for 24 hours and cell viability was established using an MTT (3-(4,5-Dimethyl-2thiazolyl)-2,5-diphenyl-2H-tetrazolium bromide) assay[22].

Data were subjected to a one-tailed t-test at (n – 2) degrees of freedom and P = 0.05.

Results and Discussion 3.1. Materials characterisation The powder XRD pattern of Ag-tobermorite is shown in Figure 1a and closely resembles those of other tobermorites reported in the literature[18,24]. Minor traces of calcite, CaCO3, (denoted by asterisks) are also present and commonly arise from atmospheric carbonation during the hydrothermal preparation of tobermorites. These data confirm that silver ions were exchanged for calcium and sodium ions within the tobermorite phase without any notable structural disruption of the lattice and that no silver-bearing precipitates were formed.

2.2. In vitro bioactivity analysis of CTAg membrane Simulated body fluid (SBF) was prepared according to the method described by Kokubo and Takadama[23]. 16 cm3 sections of CTAg membrane were placed in 20 cm3 of SBF solution in sealed polypropylene containers at 37 °C for periods of 3, 7 and 14 days. The pH of the SBF solution was measured using a Jenway 3150 pH meter and the concentrations of calcium, phosphorus and silicon were monitored by inductively coupled plasma analysis using a Perkin Elmer Optima 4300 DV spectrophotometer and multi-element standards. The recovered CTAg membranes were rinsed with deionised water, dried in air at 37 °C for 24 h and analysed by FTIR. Each analysis was carried out in triplicate.

2.3. Kirby-Bauer inhibition zone assay of CTAg and CT membranes The antimicrobial properties of the CTAg and CT membranes were assessed using the Kirby-Bauer inhibition zone method against Staphylococcus aureus NCIMB 9518, Escherichia coli NCIMB 9132 and Pseudomonas aeruginosa NCIMB 8628. Overnight cultures of each bacterium were spread on nutrient agar plates. Individual 8 mm discs of the 312

Figure 1. (a) XRD pattern of silver-exchanged tobermorite, (b) SEM image of CTAg composite membrane.

Polímeros , 25(3), 311-316, 2015

Bioactivity, biocompatibility and antimicrobial properties of a Chitosan-Mineral Composite for periodontal tissue regeneration An SEM micrograph of the CTAg composite is presented in Figure 1b and shows that the membrane’s topography is highly textured on a sub-micron scale and that the surface is characterised by fibrous tobermorite particles dispersed throughout a network of felted chitosan strands.

The FTIR spectrum of Ag-tobermorite is presented in Figure 2a. The bands at ca. 960 cm-1 and 675 cm-1 arise from various Si-O stretching modes and Si-O-Si bending vibrations, respectively[16]. O-H vibrations of silanol bonds and interlayer water molecules give rise to the broad bands at 1630 and 3450 cm–1.

The FTIR spectrum of chitosan is shown in Figure 2b.

The broad signal at ~3460 cm–1 is attributed to N-H and O-H stretching modes which overlap in this region; and the bands at 1650 and 1570 cm–1 arise from amide I C=O stretching and amide II N-H bending vibrations, respectively[25]. C-H stretching vibrations occur at 2960-2965 cm–1 and C-H bending modes give rise to the bands at 1420-1430 cm–1 and 1365 cm–1. Various C-O-C stretching frequencies occur in the range 1160 – 1060 cm-1 and the band at 1295 cm–1 is attributed to C-O-H stretching vibrations. The FTIR spectrum of the solvent-cast CTAg composite membrane, shown in Figure 2c, is essentially the sum of the individual spectra of Ag-tobermorite and chitosan with no significant shifts in the characteristic bands of either material. However, the bands at 1365, 1295 and 1160 cm-1 (which arise from aliphatic C-H, C-O-H and C-O-C groups, respectively) are less distinct and appear as poorly resolved shoulders in the spectrum of the composite owing to the overlap of the Si-O band of the tobermorite. The diminished intensity of the C-O-H band could also indicate an interaction between this group and the tobermorite lattice.

Figure 2. FTIR spectra of (a) Ag-tobermorite, (b) chitosan, (c) CTAg membrane and CTAg membrane after immersion in SBF for (d) 3 days, (e) 7 days and (f) 14 days.

Polímeros, 25(3), 311-316, 2015

3.2. In vitro bioactivity of the CTAg membrane The in vitro formation of a layer of substituted hydroxyapatite, Ca10(PO4)6(OH)2, (HA) on the surface of a material placed in SBF solution provides an indication of its bioactivity (i.e. the ability of the material to bond with living bone tissue)[23].

The formation of an HA layer on the surface of the CTAg membrane was monitored by FTIR spectroscopy and ICP analysis. The FTIR spectra of CTAg following exposure to SBF for 3, 7 and 14 days are shown in Figures 2(d), (e) and (f), respectively, and the corresponding concentrations of calcium, phosphorus and silicon species in the solution are presented in Figure 3.

Characteristic phosphate P-O bending modes of HA are noted at 570-610 cm–1 in the FTIR spectrum of the CTAg membrane after 3 days’ residence in SBF (Figure 2d).

The additional broadening of the combination band at ca.

1075 cm–1 arises from the contribution of P-O stretching modes in the 1000-1220 cm–1 range and is further evidence for the formation of HA[20]. The corresponding removal of phosphate ions and release of silicate and calcium ions are plotted in Figure 3. These data indicate that the SBF solution is essentially depleted of phosphate ions within 14 days as the HA precipitation process nears completion.

The silicate ion concentration of SBF rises from zero to 66 ppm within the first 3 days which corresponds with the dissolution of approximately 43% of the Ag-tobermorite lattice. The rate of dissolution of the Ag-tobermorite lattice then slows significantly, which is likely to be attributed to the solubility limit of silicate species in this closed system.

The degradation of the Ag-tobermorite matrix also causes a steady increase in the supernatant concentration of calcium ions and a comparatively slow release of silver ions. Silver is only detected in solution at 7 and 14 days at concentrations of 0.5 and 4.6 ppm, respectively, which correspond to the respective release of 0.6 and 1.1% of the total mass of silver present in the membrane. The delayed release of silver ions into the SBF solution during the degradation of the Ag‑tobermorite lattice may arise from ion exchange interactions with the functional groups of chitosan.

In spite of the steady release of calcium ions throughout the investigation, the pH of the SBF solution varied little about the original value of 7.45 (Figure 3). This indicates that the alkaline dissolution products of the Ag-tobermorite

Figure 3. Concentrations of P, Ca, Si and Ag in SBF and corresponding pH as functions of residence time.

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Hurt, A. P., Kotha, A. K., Trivedi, V., & Coleman, N. J. lattice are buffered by the acidic degradation products of the chitosan.

3.3. Antibacterial properties of CTAg and CT membranes The results of the antibacterial inhibition zone assays using S. aureus, P. aeruginosa and E. coli are listed in Table 1.

Distinct clear zones were noted around the silver‑bearing CTAg composite membranes in contact with all three microorganisms and in each case the bacteria were observed to readily colonise the surfaces of the control CT discs.

The microbiota of the dental biofilm in periodontitis is highly complex. In addition to the characteristic anaerobic Gram-negative oral microorganisms, the ‘non-oral’ pathogenic bacteria S. aureus, P. aeruginosa and E. coli are commonly detected within the subgingival biofilm of patients with chronic and aggressive periodontitis[26,27]. These bacteria are highly prevalent in bone and dental implant-centred infections and are among the principal pathogens associated with osteomyelitis[28,29].

The eradication of S. aureus, P. aeruginosa and E. coli in the presence of implanted biomaterials is particularly challenging owing to their biofilm formation and superior antimicrobial resistance[29]. Once bacterial adhesion has occurred on a biomaterial implant surface, the microorganisms secrete an exopolysaccharide layer that protects them from the host’s immune response and from subsequent antibiotic therapy. The host tissue cells are then unlikely to be able to displace the persistent biofilm and the implant will require surgical revision.

One strategy to promote host tissue integration over bacterial biofilm formation on implantable biomaterials is the incorporation of antibiotic components[29]. In this respect, silver salts, nanoparticles and complexes have been incorporated into a range of wound-dressings and implantable biomedical devices to exploit the broad-spectrum antimicrobial properties of this metal[18].

This research has indicated that the silver-free CT membranes readily supported colonies of S. aureus, P. aeruginosa and E. coli at concentrations of 1.0 x 109, 5.9 x 109, and 5.9 x 108 colony forming units per plate, respectively. Conversely, marked inhibitory effects were noted for each microorganism in the presence of the silver‑bearing CTAg membranes. Similar prophylaxis against microbial colonisation by S. aureus and E. coli has also been reported for composites comprising chitosan and a commercial silver-exchanged zeolite (Ag-Ion) which are intended as antimicrobial food packaging materials[30].

3.4. In vitro biocompatibility of CTAg membrane It is essential that a bioactive GTR membrane is biocompatible with bone tissue. In this respect, human osteosarcoma cells provide a model for the initial in vitro cytotoxicity appraisal of candidate biomaterials for bone tissue regeneration. The indirect cytotoxicity of the CTAg membrane towards MG63 human osteosarcoma cells was evaluated using an MTT assay. This method is derived from the ability of viable cells to reduce the tetrazolium salt of MTT into formazan which is then monitored colorimetrically.

The cell viability data for cultures in contact with increasing quantities of CTAg membrane are compared with those of the control (which consisted of cells and media only) in Figure 4. These data indicate that there is no significant loss of cell viability for the cultures in contact with 1 and 4 CTAg membrane strips compared with the control (P = 0.05).

In a similar study, the in vitro viability of human osteosarcoma cells was maintained when exposed to a porous chitosan-hydroxyapatite tissue scaffold embedded with silver nanoparticles[16]. Conversely, marked cytotoxic effects on MG63 osteosarcoma cells were observed for a composite coating comprising chitosan, Bioglass and silver nanoparticles[31]. Other research has indicated that the dosage and chemical form in which silver is presented (i.e. as ions, complexes, elemental powder or nanoparticles) impacts upon its biocompatibility with respect to mammalian cells[32]. In this case, the dosage and form of silver within the CTAg composite membrane is able to confer antimicrobial activity without compromising in vitro biocompatibility with respect to osteosarcoma cells.

3.5. CTAg as a GTR membrane Extensive clinical research has indicated that the GTR membrane-assisted regeneration of damaged periodontal tissues is favoured over traditional treatment approaches for periodontitis. The principal functions of the GTR membrane are to maintain an appropriate volume into which the compromised PDL and alveolar bone tissues can repopulate and to exclude invasion from the more fast-growing gingival and epithelial tissues. A further, desirable property of the GTR membrane would be an ability to stimulate and enhance the regeneration of the damaged tissues; and an additional asset would be the ability of the material to control bacterial activity at the wound site. This latter property is especially significant, as the damaged periodontal tissue is in contact with the external environment of the oral cavity and is at a high risk of infection during the healing process.

Table 1. Inhibition zone data for CTAg and CT membranes.

Bacterium

CTAg

CT control

S. aureus Zone of inhibition (mm)

0.81 ± 0.13

0

P. aeruginosa Zone of inhibition (mm)

1.44 ± 0.12

0

E. coli Zone of inhibition (mm)

1.07 ± 0.13

0

314

Figure 4. Viability of MG63 human osteosarcoma cells in contact with CTAg composite membrane strips.

Polímeros , 25(3), 311-316, 2015

Bioactivity, biocompatibility and antimicrobial properties of a Chitosan-Mineral Composite for periodontal tissue regeneration The principal disadvantages of chitosan in GTR applications are that its bioactivity and antimicrobial properties are insufficient and that its acidic dissolution products stimulate inflammatory reactions that inhibit healing[1,25]. This study confirms that the CTAg membrane exhibits in vitro bioactivity and that the tobermorite lattice degrades on contact with SBF to release calcium, silicate, silver and hydroxide ions. The basic calcium and hydroxide ions buffer the acidic breakdown products of the chitosan and reduce deviations from the normal physiological pH of ~7.45. Also, the dissolved silicate and calcium ions are both potent chemical signals for osteoblast activity and directed bone cell growth[1,33,34].

In addition to the chemical environment, local topography also regulates the adhesion, expression and growth of osteocytes. Mineralised alveolar bone, cementum and bone ECM possess micron and sub-micron scale textures that favour the attachment and proliferation of bone tissue cells.

In this respect, the sub-micron scale roughness of the CTAg membrane and the release of calcium and silicate ions from the tobermorite lattice are all anticipated to contribute to its therapeutic potential with respect to the regeneration of compromised alveolar bone.

The incorporation of silver cations in the composite membrane contributes significantly to its antimicrobial activity against both Gram-positive (S. aureus) and Gramnegative (P. aeruginosa and E. coli) bacteria with no observed compromise in cytocompatibility. The very slow release of silver ions from the CTAg membrane is also considered advantageous, as this is likely to provide sustained microbial resistance to bacterial biofilm formation.

Conclusions A textured polymer-mineral composite membrane has been prepared by solvent casting a mixture of chitosan and silver-bearing tobermorite. The composite membrane exhibited in vitro bioactivity in simulated body fluid and its cytocompatibility was confirmed using MG63 human osteosarcoma cells via MTT assay. The acidic degradation products of the chitosan polymer were buffered by the alkaline dissolution products of the tobermorite lattice and the slow release of sliver ions significantly enhanced the antimicrobial activity of the membrane against S. aureus, P. aeruginosa and E. coli. The results of this investigation have indicated that the chitosan-silver-tobermorite composite is a prospective candidate for GTR applications.

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16. Saravanan, S., Nethala, S., Pattnaik, S., Tripathi, A., Moorthi, A., & Selvamurugan, N. (2011). Preparation, characterization and antimicrobial activity of a bio-composite scaffold containing chitosan/nano-hydroxyapatite/nano-silver for bone tissue engineering. International Journal of Biological Macromolecules, 49(2), 188-193. http://dx.doi.org/10.1016/j. ijbiomac.2011.04.010. PMid:21549747.

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Acta Biomaterialia, 3(2), 271-276. http://dx.doi.org/10.1016/j. actbio.2006.11.003. PMid:17234465.

18. Coleman, N. J. (2009). Aspects of the in vitro bioactivity and antimicrobial properties of Ag(+)- and Zn (2+)-exchanged 11 A tobermorites. Journal of Materials Science. Materials in Medicine, 20(6), 1347-1355. http://dx.doi.org/10.1007/ s10856-009-3698-0. PMid:19214715.

19. Baek, S.-H., Plenk, H., Jr., & Kim, S. (2005). Periapical tissue responses and cementum regeneration with amalgam, SuperEBA, and MTA as root-end filling materials. Journal of Endodontics, 31(6), 444-449. http://dx.doi.org/10.1097/01. don.0000148145.81366.a5. PMid:15917684.

20. Katsamakis, S., Slot, D. E., Van der Sluis, L. W. M., & Van der Weijden, F. (2013). Histological responses of the periodontium to MTA: a systematic review. Journal of Clinical Periodontology, 40(4), 334-344. http://dx.doi.org/10.1111/ jcpe.12058. PMid:23405962.

21. Darvell, B. W., & Wu, R. C. T. (2011). “MTA”-an Hydraulic Silicate Cement: review update and setting reaction. Dental Materials, 27(5), 407-422. http://dx.doi.org/10.1016/j. dental.2011.02.001. PMid:21353694.

22. Hurt, A. P., Getti, G., & Coleman, N. J. (2014). Bioactivity and biocompatibility of a chitosan-tobermorite composite membrane for guided tissue regeneration. International Journal of Biological Macromolecules, 64, 11-16. http://dx.doi. org/10.1016/j.ijbiomac.2013.11.020. PMid:24296410.

23. Kokubo, T., & Takadama, H. (2006). How useful is SBF in predicting in vivo bone bioactivity? Biomaterials, 27(15), 29072915. http://dx.doi.org/10.1016/j.biomaterials.2006.01.017.

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24. Grangeon, S., Claret, F., Linard, Y., & Chiaberge, C. (2013).

X-ray diffraction: a powerful tool to probe and understand the structure of nanocrystalline calcium silicate hydrates.

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25. Lee, E.-J., Shin, D.-S., Kim, H.-E., Kim, H.-W., Koh, Y.-H., & Jang, J.-H. (2009). Membrane of hybrid chitosan-silica xerogel for guided bone regeneration. Biomaterials, 30(5), 743-750. http://dx.doi.org/10.1016/j.biomaterials.2008.10.025.

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26. Souto, R., Andrade, A. F. B., Uzeda, M., & Colombo, A. P. V. (2006). Prevalence of “non-oral” pathogenic bacteria in subgingival biofilm of subjects with chronic periodontitis.

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28. Pye, A. D., Lockhart, D. E. A., Dawson, M. P., Murray, C. A., & Smith, A. J. (2009). A review of dental implants and infection. The Journal of Hospital Infection, 72(2), 104-110. http://dx.doi.org/10.1016/j.jhin.2009.02.010. PMid:19329223.

29. Ribeiro, M., Monteiro, F. J., & Ferraz, M. P. (2012). Infection of orthopedic implants with emphasis on bacterial adhesion process and techniques used in studying bacterial-material interactions. Biomatter, 2(4), 176-194. http://dx.doi.org/10.4161/ biom.22905. PMid:23507884.

30. Rhim, J.-W., Hong, S.-I., Park, H.-M., & Ng, P. K. W. (2006). Preparation and characterization of chitosan-based nanocomposite films with antimicrobial activity. Journal of Agricultural and Food Chemistry, 54(16), 5814-5822. http:// dx.doi.org/10.1021/jf060658h. PMid:16881682.

31. Pishbin, F., Mouriño, V., Gilchrist, J. B., McComb, D. W., Kreppel, S., Salih, V., Ryan, M. P., & Boccaccini, A. R. (2013).

Single-step electrochemical deposition of antimicrobial orthopaedic coatings based on a bioactive glass/chitosan/ nano-silver composite system. Acta Biomaterialia, 9(7), 7469-7479. http://dx.doi.org/10.1016/j.actbio.2013.03.006.

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32. Hardes, J., Streitburger, A., Ahrens, H., Nusselt, T., Gebert, C., Winkelmann, W., Battmann, A., & Gosheger, G. (2007).

The influence of elementary silver versus titanium on osteoblasts behaviour in vitro using human osteosarcoma cell lines. Sarcoma. 2007 (Article ID 26539), 1-5. http://dx.doi. org/10.1155/2007/26539.

33. Han, P. P., Wu, C. T., & Xiao, Y. (2013). The effect of silicate ions on proliferation, osteogenic differentiation and cell signalling pathways (WNT and SHH) of bone marrow stromal cells. Biomaterials Science, 1(4), 379-392. http://dx.doi. org/10.1039/C2BM00108J. 34. Arcos, D., & Vallet-Regí, M. (2010). Sol-gel silica-based biomaterials and bone tissue regeneration. Acta Biomaterialia, 6(8), 2874-2888. http://dx.doi.org/10.1016/j.actbio.2010.02.012.

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Received: July 11, 2014 Revised: Jan. 17, 2015 Accepted: Jan. 23, 2015

Polímeros , 25(3), 311-316, 2015

http://dx.doi.org/10.1590/0104-1428.1428

Desenvolvimento da metodologia para síntese do poli(ácido lático-co-ácido glicólico) para utilização na produção de fontes radioativas Development of methodology for the synthesis of poly(lactic acid-co-glycolic acid) for use in the production of radioactive sources Fernando dos Santos Peleias Junior1,2, Carlos Alberto Zeituni1, Maria Elisa Chuery Martins Rostelato1, Guilhermino José Macêdo Fechine3, Carla Daruich de Souza1, Fábio Rodrigues de Mattos1, Eduardo Santana de Moura1, João Augusto Moura1, Marcos Antônio Gimenes Benega1, Anselmo Feher1, Osvaldo Luiz da Costa1 e Bruna Teiga Rodrigues1 Centro de Tecnologia das Radiações – CTR, Instituto de Pesquisas Energéticas e Nucleares – IPEN, CEP 05508-000, São Paulo, SP, Brasil 2 Department of Chemistry, Centre for Sustainable Chemical Technologies, University of Bath, Bath BA2 7AY, Bath, Somerset, United Kingdom 3 Escola de Engenharia, Universidade Presbiteriana Mackenzie, CEP 01302-907, São Paulo, SP, Brasil 1

*fernandopeleias@gmail.com

Resumo A Organização Mundial da Saúde (OMS) relata o câncer como uma das principais causas de morte no mundo.

Uma modalidade de tratamento que vem sendo bastante utilizada no tratamento do câncer de próstata é a braquiterapia, que consiste na introdução de sementes com material radioativo no interior do orgão. Sementes de Iodo-125 podem ser inseridas soltas ou em cordas poliméricas fabricadas a partir do (poli(ácido lático-co-ácido glicólico)) (PLGA).

Foi proposto neste trabalho, o estudo e desenvolvimento da metodologia de síntese do biopolímero PLGA. Os resultados obtidos demonstram que, através da metodologia utilizada, foi possível determinar os melhores parâmetros de reação (tempo e temperatura) para o PLGA na proporção 80/20 (lactídeo/glicolídeo). Com uma temperatura de 110 °C e tempo de reação 72h o rendimento da reação é superior a 90%. Os valores de massas moleculares obtidas entre os testes, ainda são baixos quando comparados com os valores obtidos por outros autores na literatura. Novos testes estão sendo conduzidos, utilizando dímeros preparados no laboratório. Testes substituindo o vácuo por uma atmosfera de nitrogênio também estão sendo realizados. Essas duas substituições podem aumentar o valor final da massa molecular do polímero.

Em relação à caracterização, as técnicas utilizadas confirmaram a estrutura esperada do polímero.

Palavras-chave: câncer de próstata, braquiterapia, poli(ácido lático-co-ácido glicólico), PLGA. Abstract According to the World Health Organization, cancer is a leading cause of death worldwide. A radiotherapy method extensively used in prostate cancer is brachytherapy, where the area requiring treatment receives radioactive seeds.

Iodine-125 seeds can be inserted loose or stranded in bioabsorbable polymers produced from poly(lactic-co-glycolic acid) (PLGA). We developed the synthesis methodology for PLGA and the results obtained show that it was possible to determine the optimal reaction parameters (time and temperature) for PLGA in 80/20 (lactide/glycolide) ratio. The yield was higher than 90% using a temperature of 110 °C and reaction time of 72 hours; however, the molecular weight values​​ obtained are very low compared to those obtained by other authors. New tests using previously synthesized dimers and nitrogen atmosphere are being performed. These conditions could potentially increase the molar mass of PLGA. All techniques used confirmed the expected structure of the polymer.

Keywords: prostate cancer, brachytherapy, poly(lactic-co-glycolic acid), PLGA.

Polímeros, 25(3), 317-325, 2015

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T T T T T T T T T T T T T T T T T T

Peleias, F. S., Jr., Zeituni, C. A., Rostelato, M. E. C. M., Fechine, G. J. M., Souza, C. D., Mattos, F. R., Moura, E. S., Moura, J. A., Benega, M. A. G., Feher, A., Costa, O. L., & Rodrigues, B. T.

1. Introdução A Organização Mundial da Saúde (OMS) relata o câncer como uma das principais causas de morte no mundo[1-3].

O câncer de próstata é o segundo tipo de câncer mais prevalente em homens[4,5]. Com o crescimento da expectativa de vida mundial, é esperado que o número de casos novos aumente cerca de 60% até o ano de 2015. Uma modalidade que vem sendo bastante utilizada com sucesso é a braquiterapia, que consiste na introdução de sementes de Iodo-125 radioativo no interior da próstata, próximas ao tumor, afetando ao mínimo outros órgãos nas proximidades[6,7]. No cenário brasileiro, a maior parte das sementes de Iodo-125 utilizadas são recobertas por um material polimérico bioabsorvível (poli(ácido lactico‑co‑glicólico) – PLGA). A função do PLGA é facilitar a fixação das sementes na próstata, reduzindo a ocorrência de migrações de sementes para outros órgãos[8].

O Instituto de Pesquisas Energéticas e Nucleares (IPEN) estabeleceu um projeto para fabricação de sementes de Iodo-125 e, devido às exigências do mercado nacional, surgiu a necessidade de produção de PLGA para a cobertura das sementes.

1.1. Métodos para síntese de PLGA Os processos de polimerização dos poli(α–hidroxiácidos carboxílicos) necessitam que os monômeros utilizados tenham alto teor de pureza, uma vez que as impurezas interferem nos caminhos da reação gerando polímeros de baixa qualidade. Dentre essas impurezas destacam-se a água e compostos contendo os grupos funcionais hidroxila e carboxila, que podem reduzir a massa molecular final e a taxa de polimerização, respectivamente[9]. Existem quatro métodos possíveis para produção dos poli(α–hidroxiácidos carboxílicos): polimerização por policondensação direta, policondensação azeotrópica, polimerização em estado sólido (SSP), e polimerização por abertura de anéis (ROP)[9,10].

Nos dois primeiros métodos há eliminação de moléculas de água e, de um modo geral, o polímero final obtido possui baixa massa molecular. Embora a policondensação azeotrópica

permita a obtenção de massas moleculares maiores, existe o problema da toxicidade do solvente associada[9,11]. Através da polimerização no estado sólido, também é possível preparar polímeros com massas moleculares elevadas.

No entanto, o método exige altos tempos de reação, e também há necessidade de um pré-polímero de baixa massa molecular[9,12,13]. Por último, a polimerização por abertura de anel se destaca como método preferido para produção de PLGA. Neste tipo de polimerização, parte-se dos dímeros cíclicos lactídeo e glicolídeo. Através da abertura dos anéis dos dímeros, tem-se a geração de uma bifuncionalidade, que ao reagir com outras unidades monoméricas, forma uma cadeia polimérica. Neste tipo de polimerização não há a formação de subprodutos durante a reação. Os dímeros cíclicos são produzidos a partir da despolimerização de polímeros de baixa massa molecular obtidos por policondensação[9,11,14,15].

A síntese de polímeros por abertura de anel é o método mais comumente estudado devido a possibilidade de variar as propriedades dos polímeros resultantes de forma mais controlada, ampliando o seu campo de aplicação. Como os dímeros lactídeo e glicolídeo podem apresentar impurezas, como a água, ou até mesmo os monômeros que não reagiram, é de extrema importância que um processo cuidadoso de purificação desses dímeros seja realizado[9].

O mecanismo mais utilizado para esse tipo de polimerização é o de coordenação-inserção, pois permite a obtenção de materiais com alta massa molecular[9]. A Figura 1 ilustra o mecanismo de coordenação inserção, através da polimerização do lactídeo, utilizando alcóxido de alumínio como catalisador.

O primeiro passo do mecanismo ocorre quando o oxigênio exocíclico do dímero forma temporarioamente um composto de coordenação com o metal presente no catalisador.

Essa ligação aumenta a nucleofilicidade do catalisador (no exemplo acima, o alcóxido de alumínio) , assim como a eletrofilicidade do grupo carbonila presente no dímero.

No segundo passo, a ligação acila – oxigênio é quebrada e o lactídeo já aberto, é inserido no ligação metal-oxigênio do catalisador[9,16,17]. A reação continua conforme os anéis

Figura 1. Demonstração do mecanismo de coordenação-inserção na polimerização em PLA[9].

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Polímeros , 25(3), 317-325, 2015

Desenvolvimento da metodologia para síntese do poli(ácido lático-co-ácido glicólico) para utilização na produção de fontes radioativas de outras moléculas de lactídeo são abertos e inseridos no catalisador entre o átomo do metal e oxigênio adjacente[16,17].

Variando-se condições da reação (tempo, temperatura e concentração do catalisador), polímeros de diferentes massas moleculares são obtidos. O mecanismo de reação do glicolídeo é o mesmo, pois a única diferença é a ausência dos radicais metila. Embora uma grande variedade de catalisadores tenha sido pesquisada, o 2-etilhexanoato de estanho (octanoato de estanho – Sn(Oct)2 é o mais utilizado em aplicações biomédicas devido a sua alta eficiência e baixa toxicidade[9,18-20].

As vantagens da polimerização por abertura de anel são: menor tempo de reação e maiores taxas de conversão do monômero. Como desvantagens destacam-se: custos mais altos de produção e dificuldades técnicas para obtenção dos dímeros cíclicos com elevada pureza[18,19].

O objetivo deste trabalho é desenvolver uma possível metodologia de síntese PLGA, utilizando o método de polimerização por abertura de anéis, com o propósito de utilizar o material sintetizado para recobrir as sementes de Iodo-125 fabricadas no laboratório do IPEN.

2. Experimental Nas reações propostas neste trabalho, foram utilizados os dímeros L-lactídeo e glicolídeo, já purificados, na proporção de 80:20 respectivamente. O motivo da escolha desta proporção se deu pelo fato do glicolide ser extremamente insolúvel em quase todos os solventes orgânicos, com exceção dos solventes fluorados. No entanto, a utilização de destes solventes tornaria o projeto inviável, uma vez que 100mL de hexafluoroisopropanol na empresa Sigma‑Aldrich, tem um custo aproximado de R$ 2000,00. Não sendo possível sintetizar um produto similar ao Vycril (90% Glicolídeo/ 10% Lactídeo), que já é utilizado em braquiterapia prostática, optou-se

pela utilização da proporção descrita acima, que também é utilizada na fabricação de suturas para fixação de tendões e ligamentos aos ossos[21,22].

Uma vez determinada a proporção dos monômeros desejada, determinou-se quais parâmetros seriam avaliados nas reações. Considerando que a reação será conduzida em bulk, ou seja, sem adição de solventes, os principais parâmetros a serem avaliados seriam: tempo, temperatura e concentração do catalisador. Neste trabalho optou-se por avaliar apenas a influência da temperatura e do tempo de reação, pois a concentração do catalisador (SnOct2) de 1/5000 é unânime em todas as referências consultadas. Em relação ao tempo de reação e temperatura, existe uma grande faixa de valores utilizados, que varia de autor para autor.

2.1. Síntese do poli(ácido lático-co-ácido glicólico) Os dímeros L-lactídeo e glicolídeo, e o catalisador octanoato de estanho (SnOct2), utilizados nas reações, foram adquiridos da empresa Sigma-Aldrich. As massas correspondentes dos dímeros e do catalisador foram adicionadas a uma ampola de vidro, que foi posteriormente selada a vácuo, a fim de que a reação se desse na ausência de O2 e possível umidade contida no ar.

As ampolas foram imersas em um banho termostático da marca Lauda, modelo E200. Foram escolhidas 3 faixas de temperatura (110 °C, 140 °C, 170 °C) e 4 diferentes tempos de reação (6h, 24h, 72h, 168h). Esses valores foram escolhidos baseados nos diferentes parâmetros encontrados nas referências estudadas[18,20,23-25]. Cada teste foi realizado em duplicata, totalizando 24 reações. As ampolas foram numeradas de 1 a 24. A Tabela 1 relaciona o número de cada ampola com os parâmetros utilizados, assim como as massas dos monômeros e catalisador pesadas.

Decorrido o tempo proposto para cada reação, as ampolas foram retiradas do banho e mergulhadas em banho

Tabela 1. Valores de massa e parâmetros utilizados nas reações de polimerização.

Nº da Ampola 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24

Massa de L-Lactídeo (g) 5,973 5,972 5,973 5,971 5,982 5,979 5,974 5,977 5,970 5,973 5,972 5,979 5,976 5,973 5,980 5,979 5,970 5,970 5,971 5,970 5,971 5,971 5,971 5,972

Polímeros, 25(3), 317-325, 2015

% em mol de L-lactídeo 79,91 79,92 79,93 79,99 79,92 80,00 79,89 79,92 79,93 79,97 79,99 80,01 79,96 79,90 80,01 80,01 79,96 80,00 79,98 80,00 79,98 79,98 80,00 79,99

Massa de glicolídeo (g) 1,209 1,208 1,208 1,203 1,210 1,204 1,211 1,209 1,207 1,205 1,203 1,203 1,206 1,210 1,203 1,203 1,205 1,202 1,204 1,202 1,204 1,204 1,202 1,203

% em mol de glicolídeo 20,09 20,08 20,07 20,01 20,08 20,00 20,11 20,08 20,07 20,03 20,01 19,99 20,04 20,10 19,99 19,99 20,04 20,00 20,02 20,00 20,02 20,02 20,00 20,01

Massa de SnOct2 0,0052 0,0048 0,0054 0,0053 0,0048 0,0049 0,0056 0,0053 0,0061 0,0052 0,0049 0,0052 0,0049 0,0052 0,0053 0,0050 0,0042 0,0046 0,0039 0,0043 0,0052 0,0051 0,0055 0,0050

Tempo (h) 6 6 24 24 72 72 168 168 6 6 24 24 72 72 168 168 6 6 24 24 72 72 168 168

Temperatura (°C) 110 110 110 110 110 110 110 110 140 140 140 140 140 140 140 140 170 170 170 170 170 170 170 170

319

Peleias, F. S., Jr., Zeituni, C. A., Rostelato, M. E. C. M., Fechine, G. J. M., Souza, C. D., Mattos, F. R., Moura, E. S., Moura, J. A., Benega, M. A. G., Feher, A., Costa, O. L., & Rodrigues, B. T. de água com gelo por 15 minutos, de modo que a reação fosse interrompida. A seguir a ampola foi quebrada, e o polímero resultante foi dissolvido, sob agitação constante, em 50mL de clorofórmio (Synthlab). O tempo para dissolução completa do polímero variou de 1 a 12 horas, dependendo dos parâmetros da reação. Finalizada a dissolução, o polímero foi precipitado através do gotejamento de metanol (Synthlab), na solução que continha o polímero dissolvido. O metanol deve ser adicionado até o ponto em que a solução polímero + clorofórmio deixe de tornar-se turva com o gotejamento.

Após a precipitação, o polímero obtido foi submetido à filtração à vácuo. Foi utilizado papel de filtro grau 589/3, que possui velocidade de filtração lenta, mas é o mais eficiente na coleta de partículas pequenas (<2μm). Logo após a filtração do polímero, o material foi submetido ao processo de secagem, sendo primeiramente levado a uma estufa por 1 hora a uma temperatura de 80 °C. Posteriormente, o polímero foi levado a um dessecador à vácuo por 24 horas, de modo que qualquer resíduo de solvente ainda presente fosse eliminado. Finalmente, o material seco foi pesado e o rendimento pode ser calculado.

2.2. Determinação da massa molecular por cromatografia de permeação em gel (GPC) Para a obtenção dos dados referentes ás massas moleculares das 24 amostras de polímeros produzidos, utilizou-se um cromatógrafo HT-GPC-module 350A da marca Viscotek, equipado com colunas GPC HT-806M da marca Shodex. Em pequenos frascos, as amostras dos polímeros foram dissolvidas em tetrahidrofurano (THF) e posteriormente filtradas em filtros Millipore com abertura de 0,45μm. As condições de operação utilizadas foram as seguintes: taxa de bombeamento de 1ml/min, temperatura

de 40 °C (recomendada quando utilizado o solvente THF) e volume de injeção de 100μL[18,20,23,26].

2.3. Espectroscopia na região do infravermelho (IV) Os espectros das 24 amostras foram obtidos em um espectrômetro FT-IR da marca Perkin Elmer, modelo Spectrum 100. O espectrômetro estava equipado com um acessório de refletância total atenuada (ATR), com cristal de diamante. Este acessório elimina o tradicional método de preparação de amostras pela técnica da pastilha de brometo de potássio, tornando a análise ainda mais rápida[27].

Os espectros foram obtidos a temperatura ambiente, no estado sólido, utilizando o intervalo de 650-4000 cm–1.

A resolução utilizada foi de 4 cm–1.

2.4. Espectroscopia Raman Os espectros foram obtidos em um espectrômetro Raman Horiba – Jobin Yvon, modelo XploRA. Não é necessário preparação das amostras, e os espectros foram obtidos em condições ambientais normais. A fonte de luz visível tinha comprimento de onda de 532 nm e a potência do laser foi ajustada para 50%. Os espectros foram obtidos no intervalo entre 350-3500 cm–1 com resolução de 1,8 cm–1. Os parâmetros utilizados foram sugeridos pelo fabricante.

3. Resultados e Discussão 3.1. Resultados dos rendimentos e aspectos gerais do material produzido Nesta seção serão apresentados os rendimentos obtidos nas reações (Tabela 2) e os aspectos gerais dos materiais produzidos em diferentes condições.

Tabela 2. Valores de massa e rendimento obtidos nas reações de polimerização.

Massa dos Tempo (h) Reagentes (g) 1 7,187 6 2 7,185 6 3 7,186 24 4 7,179 24 5 7,197 72 6 7,188 72 7 7,191 168 8 7,191 168 9 7,183 6 10 7,183 6 11 7,180 24 12 7,187 24 13 7,187 72 14 7,188 72 15 7,188 168 16 7,187 168 17 7,179 6 18 7,177 6 19 7,179 24 20 7,176 24 21 7,180 72 22 7,180 72 23 7,179 168 24 7,180 168 * Não houve precipitação após adição de metanol.

Nº da Ampola

320

Temperatura (°C) 110 110 110 110 110 110 110 110 140 140 140 140 140 140 140 140 170 170 170 170 170 170 170 170

Massa obtida após purificação (g) 4,381 4,035 6,378 6,103 6,797 6,563 5,644 5,871 4,890 5,080 5,340 5,620 3,674 2,954 0,222 0,287 2,554 2,623 2,143 1,916 0,763 0,949 * *

Rendimento (%) 60,96 56,16 88,76 85,01 94,44 91,30 78,49 81,64 68,07 70,72 74,38 78,20 51,18 41,10 3,08 3,99 35,58 36,55 29,85 26,69 10,63 13,22 * *

Média do Rendimento 58,56 86,88 92,87 80,06 69,40 76,29 46,14 3,54 36,06 28,27 11,92 *

Polímeros , 25(3), 317-325, 2015

Desenvolvimento da metodologia para síntese do poli(ácido lático-co-ácido glicólico) para utilização na produção de fontes radioativas Analisando a Tabela 2, é possível concluir que, considerando apenas o rendimento, os parâmetros indicados para uma produção em série seriam 72 horas e 110 °C, correspondentes às ampolas 5 e 6. Utilizando estes valores, o rendimento foi superior a 90%. É possível concluir também que a reação não se processa adequadamente em temperaturas elevadas, visto que o máximo de rendimento obtido a 170 °C foi de 36,06%. O tempo de reação está diretamente ligado à temperatura utilizada. Em temperaturas menores, melhores resultados são obtidos em tempos maiores (porém, tempos demasiadamente longos podem ocasionar uma diminuição no rendimento – vide reações 7 e 8). O inverso ocorre em altas temperaturas, uma vez que há uma diminuição contínua do rendimento com o aumento do tempo de reação.

Embora ótimos rendimentos sejam altamente desejáveis, outros parâmetros e condições do material produzido são igualmente importantes. A seguir são feitos alguns comentários observados ao longo do procedimento utilizado.

Nas ampolas 1 e 2 (110°C-6h), o aspecto do polímero após o término da reação era de um líquido bem viscoso e transparente. Este aspecto se manteve mesmo após o resfriamento. Considerando o ponto de fusão do PLGA, na proporção 80/20, em torno de 140 °C, e dos monômeros em torno de 100ºC, conclui-se que o grau de polimerização dos monômeros foi baixo. Após o procedimento de purificação, o aspecto pegajoso se manteve, sendo muito difícil de retirar completamente a amostra do béquer. A filtração de ambas as amostras foi rápida – cerca de 15 minutos.

Nas ampolas 3 e 4 (110 °C-24h), ao término da reação, o material estava no estado sólido e tinha aspecto extremamente duro e levemente esbranquiçado. O tempo de dissolução em clorofórmio foi muito maior que os testes 1 e 2 (cerca de 12 horas), sugerindo massas moleculares maiores. Após a precipitação em metanol, as amostras também apresentaram aspecto pegajoso. Entretanto, ao tirar as amostras do solvente, em poucos segundos o material perdia a característica pegajosa e passava a se assemelhar a uma cola depois de seca. O processo de filtração foi semelhante ao dos testes 1 e 2, levando cerca de 15 minutos para completar a operação. Um detalhe importante observado nos testes 1-4, é que o polímero não desgrudou da ampola após o término da reação. Foi necessário adicionar a amostra + ampola no clorofórmio, com posterior remoção da ampola após completa dissolução. Após a purificação, o aspecto do material era de um pó branco ligeiramente empedrado. O mesmo aspecto também foi observado nas reações das ampolas 5-20.

O material das ampolas 5 e 6 (110 °C-72h) e 7 e 8 (110 °C-168h) apresentaram características muito semelhantes aos anteriores (3-4). Diferenças foram observadas em relação a cor (os 4 últimos já estavam bastante esbranquiçados), e em relação à aderência do polímero na ampola. Os 4 últimos testes foram separados com muita facilidade da ampola, e foram adicionados diretamente no clorofórmio. O tempo de filtração foi ainda menor que os anteriores (cerca de 5 minutos).

Em relação aos testes da próxima bateria (140 °C), o aspecto das ampolas 9-14 era de um líquido bem viscoso e transparente, assim como nos testes 1 e 2. Entretanto, a solidificação do material ocorreu logo após a retirada das Polímeros, 25(3), 317-325, 2015

ampolas do banho termostático. Isto ocorreu pelo fato da reação ser conduzida acima do ponto de fusão do polímero na proporção 80/20. Diferenças foram encontradas na etapa de purificação do polímero. O líquido contendo clorofórmio + metanol + polímero dissolvido apresentou aspecto turvo com uma coloração esbranquiçada forte, pelo fato de conter ainda em solução o polímero/monômero não reagido.

Mesmo após passar pelo filtro, a solução continuava com o aspecto esbranquiçado. Esta condição tornou o processo de filtração muito mais demorado (>2h) e ineficiente, diminuindo o rendimento.

Com o objetivo de aumentar o rendimento destas reações, foi adicionado novamente clorofórmio no líquido já filtrado, e em seguida ele foi reprecipitado para ser filtrado pela segunda vez. O aumento no rendimento foi muito pequeno (~3%), não justificando a realização de múltiplas filtrações em todos os testes. Os dois últimos testes desta bateria saíram do banho termostático com uma coloração amarelada, e praticamente não houve precipitação da solução polímero + clorofórmio após adição do metanol, mesmo com o dobro do solvente utilizado (~500mL). A partir do teste 13, foram encontradas dificuldades na manipulação do polímero, uma vez que ele voltou a se fixar nas paredes da ampola.

Por fim, na última bateria de testes (17-24), todos os polímeros apresentaram coloração amarelada, aumentando a intensidade da cor com o aumento do tempo de reação.

A diferença na coloração sugere uma possível degradação do material devido à altas temperaturas e tempos de reação.

Em relação ao rendimento, este diminuiu conforme o tempo de reação tornava-se maior, chegando a 0% no último caso.

3.2. Resultados das análises de cromatografia de permeação em gel (GPC) Os resultados dos diferentes tipos de massas moleculares (Mn, Mw e Mz), assim como a média da Mw, a massa dos 10% das cadeias mais pesadas e o índice de polidispersividade são apresentados na Tabela 3.

Analisando a Tabela 3, é possível observar que os polímeros que apresentaram maiores valores de massas moleculares são os dos testes 5 e o 6 (com média de Mw = 17106 Daltons), seguidos pelos dos testes 3 e 4 (com média de Mw = 15393 Daltons). É possível observar também que alguns testes apresentam dois valores de massa molecular, sendo um deles indicado com um asterisco. Isso ocorreu pelo fato do surgimento de dois picos ao longo da eluição, com massas moleculares muito diferentes. Este segundo pico, que apresentou em quase todos os casos, massas moleculares muito próximas do valor da massa dos próprios dímero, pode ter sido gerado por material que não reagiu.

Mesmo considerando às baixíssimas massas moleculares em um cálculo separado, o valor do Mw obtido está abaixo do esperado quando se realiza o processo de polimerização por aberturas de anéis, partindo dos dímeros já purificados.

Este valor se encontra próximo aos 100 kDalton. Como já foi citado anteriormente, a reação é extremamente sensível a algumas impurezas, entre elas a água. Além disso, os dímeros utilizados são extremamente reativos, e embora tenham sido armazenados nas condições recomendadas pelo fabricante (freezer), uma parte pode ter sofrido hidrólise e comprometido o crescimento das cadeias. Novos testes 321

Peleias, F. S., Jr., Zeituni, C. A., Rostelato, M. E. C. M., Fechine, G. J. M., Souza, C. D., Mattos, F. R., Moura, E. S., Moura, J. A., Benega, M. A. G., Feher, A., Costa, O. L., & Rodrigues, B. T. Tabela 3. Valores das diferentes massas moleculares obtidos por GPC. Nº da Ampola 1 2 1* 2* 3 4 3* 4* 5 6 5* 7 8 8* 9 10 9* 10* 11 12 12* 13 14 15 16 17 18 19 20 21 22 21* 22* 23 24 23*

Mn

Mw

(Daltons) 4173 7322 297 233 8917 9862 231 225 9983 10771 169 2024 7169 227 4881 5351 288 236 4741 5617 244 2111 2158 1756 1300 1117 1146 4126 1131 3348 4166 230 228 2950 2198 273

(Daltons) 6517 10716 319 256 14678 16108 257 235 16285 17926 211 6306 15215 254 8151 9149 312 256 9886 11216 254 4512 4986 3819 2559 2142 2130 8598 2283 7351 9178 243 240 5788 5741 300

Média Mw (Daltons) 8616 288 15393 246 17106 211 10760 254 8650 284 10551 254 4749 3189 2136 5440 8264 242 5764 300

Mz

Mw 10% maiores

(Daltons) 8593 13950 340 281 20910 22916 283 246 23460 25892 260 12843 24583 283 12131 13859 335 278 16098 17745 264 7257 8301 6287 4135 3456 3361 14410 3991 12119 15134 257 254 9502 9669 328

(Daltons) 14119 23565 475 415 36210 39547 425 332 40607 44699 421 21563 42185 419 20987 24052 474 412 27627 30412 352 12417 14171 10737 7085 5925 5766 24773 6795 20701 25845 352 349 16256 16445 458

Mw/Mn 1,561 1,463 1,076 1,097 1,646 1,633 1,111 1,046 1,631 1,664 1,251 3,115 2,122 1,121 1,670 1,710 1,083 1,084 2,085 1,997 1,040 2,137 2,311 2,175 1,968 1,917 1,859 2,084 2,019 2,132 2,203 1,060 1,055 1,961 2,611 1,099

*Amostras contendo 2 picos diferentes.

utilizando dímeros preparados no próprio laboratório estão sendo realizados. Dessa maneira, os dímeros seriam usados muito mais rapidamente, minimizando os problemas citados acima. Existem relatos na literatura da substituição do vácuo por uma atmosfera inerte, geralmente nitrogênio. Testes adicionais utilizando um fluxo contínuo de nitrogênio em um reator químico também estão sendo realizados.

Relacionando a Tabela 3 (rendimentos) com a Tabela 4 (massas moleculares) é possível concluir que os testes 5 e 6, seriam mais indicados para uma produção em série. Ambos apresentaram os maiores valores de rendimento e massas moleculares. Os testes 3 e 4 também apresentaram bons rendimentos, boa manuseabilidade e massas moleculares próximas às maiores obtidas (5 e 6). Considerando que o tempo de reação dos testes 3 e 4 (24h) é apenas 1/3 do tempo de reação dos testes 5 e 6 (72h), ele também se mostra uma alternativa, caso seja necessário maiores quantidades do polímero em menos tempo.

322

3.3. Espectroscopia na região do infravermelho (IV) Os polímeros sintetizados foram analisados estruturalmente através de espectroscopia na região IV. Todos os 24 espectros obtidos apresentaram bandas de absorção na mesma região.

A Figura 2 apresenta os espectros dos testes 5 e 6, que são aqueles com interesse prático em uma eventual produção.

As bandas destacadas nos espectros acima são representadas na Tabela 4 com as suas respectivas atribuições.

Analisando a tabela com as respectivas bandas de absorção na região do infravermelho, observa-se que todos os grupos funcionais esperados foram de fato encontrados.

As amostras dos copolímeros submetidas à análise de espectroscopia IV não apresentaram bandas de absorção entre 3700 e 3100 cm–1, característica de grupamento OH, indicando ausência de umidade nas mesmas. A Figura 3 representa a estrutura do PLGA esperada após a síntese, facilitando a relação dos grupos mostrados na Tabela 4 com os grupos esperados.

Polímeros , 25(3), 317-325, 2015

Desenvolvimento da metodologia para síntese do poli(ácido lático-co-ácido glicólico) para utilização na produção de fontes radioativas Tabela 4. Atribuição das bandas de absorção na região do infravermelho[18,28-32].

Banda de Absorção (cm–1)

Atribuição

3000-2950 1750 1454

Estiramento antissimétrico CH3 e CH2 Estiramento C=O do COO Deformação CH3 e CH2

1380-1360

Deformação CH3 e CH

1180

Deformação C-O do COO

1130-1040

Estiramento C-O do CH-O ou CH2-O

865

Estiramento C-COO

755

Deformação CH

Por último, para fins de comparação, a Figura 4 representa o espectro do copolímero PLGA obtido na literatura.

Comparando os espectros obtidos neste trabalho com o obtido por Hummel[33], nota-se grande similaridade entre eles.

Assim, pode-se concluir que os resultados obtidos para as análises de FTIR estão de acordo com a literatura. Observase uma banda entre 3400 e 3600 cm–1 no espectro mostrado na Figura 3, característica do grupo hidroxila, ausente nos espectros encontrados neste trabalho. Essa banda descrita na literatura, talvez pode ser devido à presença de umidade na amostra, de solvente residual, ou da absorção de (OH) do grupo carboxila resultante da hidrólise do monômero[34].

3.4. Espectroscopia Raman Foram escolhidos os dois melhores testes (5-6) para análise por espectroscopia Raman. Como as duas amostras apresentaram resultados idênticos, apenas uma é apresentada na Figura 5.

Assim como na espectroscopia IV, todos os picos esperados foram encontradas. A descrição das bandas encontradas no espectro é descrita na Tabela 5.

Figura 2. Espectro na região IV do copolímero PLGA 80:20 – Testes 5 e 6.

Figura 3. Estrutura esperada após síntese do copolímero PLGA[29].

Figura 5. Espectro Raman referente ao PLGA obtido no teste 6.

Figura 4. Espectro na região IV referente ao copolímero PLGA obtido na literatura[33].

Polímeros, 25(3), 317-325, 2015

323

Peleias, F. S., Jr., Zeituni, C. A., Rostelato, M. E. C. M., Fechine, G. J. M., Souza, C. D., Mattos, F. R., Moura, E. S., Moura, J. A., Benega, M. A. G., Feher, A., Costa, O. L., & Rodrigues, B. T. Tabela 5. Atribuição das bandas de absorção em espectroscopia Raman[32,35].

Banda de Absorção (cm-1)

Atribuição

3006-2950 2880 1765

Estiramento antissimétrico CH3 e CH2 Estiramento CH Estiramento C=O

1453

Deformação antissimétrica CH3

1387

Deformação simétrica CH3

1296

Deformação CH

1130

Rotação antissimétrica CH3

1090

Estiramento simétrico COC

1042

Estiramento C-CH3

878

Estiramento C-COO

720

Deformação C=O (Lactídeo)

400

Deformação CCO

4. Conclusões Os resultados obtidos demonstram que, através da metodologia proposta para síntese dos polímeros, foi possível determinar os melhores parâmetros de reação (tempo e temperatura) para o PLGA na proporção 80/20.

Com uma temperatura de 110 °C e tempo de reação de 24h, foi possível obter 86% de rendimento, e aumento o tempo de reação para 72h, na mesma faixa de temperatura, o rendimento é superior a 90%. Os resultados obtidos por GPC, também mostram que as maiores massas moleculares foram obtidas nos testes com os maiores valores de rendimento.

No entanto, os resultados de massa molecular obtidos são baixos quando comparados com os valores obtidos por outros autores na literatura. Novos testes estão sendo conduzidos, utilizando dímeros preparados no laboratório.

Como o tempo entre o preparo e a utilização do dímero seria bem mais curto, problemas envolvendo a hidrólise do material seriam evitados. Testes substituindo o vácuo por uma atmosfera de nitrogênio também estão sendo realizados.

Essas duas substituições podem aumentar o valor final da massa molecular do polímero. Em relação à caracterização da estrutura da molécula, as técnicas de espectroscopia IV e Raman confirmaram a estrutura da molécula esperada.

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http://dx.doi.org/10.1590/0104-1428.1468

Laminados biodegradáveis de blendas de amido de C mandioca e poli(vinil álcool): efeito da formulação C sobre a cor e opacidade C Sheets of cassava starch - poly (vinyl alcohol) produced by C extrusion: Effect of formulation on the color and opacity C Juliano Zanela *, Mônica Oliveira Reis , Adriana Passos Dias , Suzana Mali , Maria Victória Eiras Grossmann e Fabio Yamashita C Departamento de Ciência e Tecnologia de Alimentos, Universidade Estadual de Londrina – UEL, C CEP 86051-990, Londrina, PR, Brasil C Resumo C Foram produzidos 12 laminados com diferentes proporções de amido de mandioca, poli(vinil álcool) (PVA) e glicerol por extrusão usando um planejamento de mistura. A opacidade dos materiais variou entre 31 e 56%, e a concentração de e a interação entre as concentrações de amido e PVA foram os principais responsáveis pelo aumento da opacidade.

C PVA A diferença de cor (ΔE*) variou entre 20-30, sendo a concentração de amido o principal responsável pelo aumento da C diferença de cor, por promover uma coloração mais amarelada nos laminados.

Palavras-chave: extrusão, biopolímeros, propriedades ópticas.

C 1

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*julianozanela@gmail.com

Abstract Twelve sheets were produced by extrusion with different proportions of cassava starch, poly (vinyl alcohol) (PVA) and glycerol using a mixture design. The opacity of the materials ranged from 31 to 56%, and the concentration of PVA and the interaction between starch and PVA concentrations the mainly responsible for increasing opacity. The color difference (ΔE*) ranged from 20 to 30, the concentration of starch being primarily responsible for increasing in color difference by promoting a more yellowish color in sheets.

Keywords: extrusion, biopolymers, optical properties.

1. Introdução Há uma crescente busca pelo desenvolvimento de polímeros biodegradáveis para a substituição dos polímeros convencionais devido aos problemas ambientais e escassez de petróleo. O amido é um biopolímero bastante estudado no desenvolvimento de materiais biodegradáveis, devido ao seu baixo custo e abundância. O amido não é um termoplástico verdadeiro, mas na presença de agentes plastificantes, calor e cisalhamento ele perde sua estrutura semicristalina, originando o amido termoplástico (ATp), que é uma material amorfo, com características semelhantes as observadas em polímeros sintéticos[1,2].

Materiais produzidos somente com ATp apresentam propriedades mecânicas e barreira ao vapor de água inadequadas para produção e aplicação em escala comercial[3]. Dessa forma, é necessário produzir blendas com outros polímeros biodegradáveis para melhorar as propriedades do material, como o poli(vinil álcool) (PVA). Dependendo do grau de hidrólisee viscosidade do PVA, é possível obter materiais com diferentes propriedades[4].

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Uma característica importante dos materiais biodegradáveis é a opacidade, principalmente para aqueles que vão ser utilizados como embalagem, pois a transparência é importante para aqueles produtos que precisam ser visualizados pelo consumidor. Já para os produtos sensíveis à fotodegradação, é importante que a embalagem seja opaca[5].

O objetivo foi estudar a influência das concentrações de amido, glicerol e PVA sobre a cor e opacidade de laminados produzidos por extrusão.

2. Experimental Foi utilizado amido de mandioca (Indemil, Brasil), glicerol (Dinâmica, Brasil) e poli(vinil álcool) Selvol 203 (Sekisui Chemical, Japão), com grau de hidrólise de 88,14% e viscosidade de 4,10cP em solução a 4%, e as formulações estão na Tabela 1. As misturas foram homogeneizadas manualmente e mantidas em estufa à vácuo (Quimis, Brasil), por uma hora a 85 °C para melhorar a incorporação do

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Laminados biodegradáveis de blendas de amido de mandioca e poli(vinil álcool): efeito da formulação sobre a cor e opacidade glicerol, segundo metodologia proposta por Jang e Lee[6].

Após esta etapa, as blendas foram processadas em extrusora dupla-rosca co-rotativa modelo D-20 (BGM, Brasil), com velocidade dos parafusos de 100 rpm e perfil de temperatura de 90/170/170/170/170°C, equipada com uma matriz do tipo “flat die” com abertura de 0,8 mm, para a produção de laminado em conjunto com uma calandra de três rolos (AX Plásticos, Brasil). As velocidades dos rolos tracionadores da calandra foram definidas individualmente para cada tratamento, para garantir um laminado contínuo, coeso e de espessura homogênea.

As análises de cor e opacidade foram realizadas conforme metodologia descrita por Sobral[7], oito replicatas de cada laminado foram analisadas utilizando um colorímetro BYK Gardner (Alemanha), com um ângulo de 10° e iluminante D65.

A opacidade (Y) foi determinada como a relação entre as leituras de opacidade das amostras sobre os padrões preto (Yp) e branco (Yb), sendo os resultados expressos em escala arbitrária (0-100%) conforme Equação 1:

Y =

(Yp / Yb ) *100 (1)

A análise de cor foi determinada como a diferença de cor (ΔE*), os parâmetros ΔL*, Δa* e Δb* foram determinados pela Equação 2, como a diferença entre as leituras do padrão branco e das amostras sobre esse padrão.

2 2 2 ∆E* = ( ∆L *) + ( ∆a *) + ( ∆b *)   

0,5

(2)

O ajuste do modelo do planejamento de mistura foi realizado utilizando o software Statistica 7.0 (Statsoft, EUA).

3. Resultados e Discussão A opacidade variou de 31 a 56%, já a diferença de cor variou entre 20 e 30 quando comparado ao padrão de cor branca (Tabela 1).

A Figura 1 apresenta as superfícies de resposta e os coeficientes do modelo, sendo possível observar que o PVA foi o componente puro de maior efeito para o modelo (38,86),

Tabela 1. Componentes reais, pseudo-componentes e respostas do planejamento de mistura.

Código T1 T2 T3 T4 T5 T6 T7 T8 T9 T 10 T 11 T 12

Amido 0,675 0,630 0,585 0,540 0,600 0,560 0,520 0,480 0,525 0,490 0,455 0,420

Componente real PVA Glicerol 0,075 0,250 0,070 0,300 0,065 0,350 0,060 0,400 0,150 0,250 0,140 0,300 0,130 0,350 0,120 0,400 0,225 0,250 0,210 0,300 0,195 0,350 0,180 0,400

Amido 0,944 0,778 0,611 0,444 0,667 0,518 0,370 0,222 0,389 0,259 0,130 0,000

Pseudo-componente PVA 0,056 0,037 0,018 0,000 0,333 0,296 0,259 0,222 0,611 0,556 0,500 0,444

Glicerol 0,000 0,185 0,370 0,556 0,000 0,185 0,370 0,556 0,000 0,185 0,370 0,556

Resposta Opacidade ΔE* 39 (±2) 30 (±3) 34 (±4) 27 (±2) 36 (±4) 27 (±4) 31 (±6) 22 (±3) 55 (±4) 28 (±5) 48 (±2) 28 (±2) 43 (±3) 22 (±2) 41 (±2) 23 (±2) 56 (±3) 25 (±3) 49 (±2) 22 (±3) 46 (±3) 21 (±2) 44 (±4) 20 (±2)

Figura 1. Superfície de resposta das variáveis opacidade e diferença de cor em função da concentração de amido, PVA e glicerol, em termos de seus pseudo-componentes. β1 = amido, β2= PVA, β3 = glicerol, β12 = interação amido + PVA, β13 = interação amido + glicerol.

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Zanela, J., Reis, M. O., Dias, A. P., Mali, S., Grossmann, M. V. E., & Yamashita, F. bem como a interação amido/PVA (78,03), apresentando ainda um bom ajuste aos dados experimentais (R2= 0,82).

Para a variável ΔE*, o amido (30,91) apresentou o maior efeito, pois laminados com maiores proporções de amido apresentaram coloração mais amarelada, enquanto que laminados com maiores concentrações de PVA apresentaram coloração mais esbranquiçada, resultando em menor ΔE* e maior opacidade.

Moraes et al.[8] e Silva et al.[9] em filmes de PVA e gelatina plastificados com glicerol e obtidos por casting obtiveram valores máximos de opacidade de 2,2 e 2,4% respectivamente, e ΔE* de no máximo 3,0 e 3,6, respectivamente, muito inferiores aos obtidos no presente estudo, isso se deve ao uso de gelatina ao invés de amido, a menor espessura dos materiais obtidos, bem como a obtenção pela metodologia de casting, que promove a formação de materiais mais uniformes, transparentes e menos espessos quando comparados aos laminados obtidos por extrusão. Devido a sua estrutura, a gelatina ao contrário do amido não apresenta cristalinidade, permitindo uma maior passagem da luz por estar em um estado totalmente amorfo[10], porém tanto o PVA quando o amido sofrem recristalização, isso reduz a passagem de luz, aumentando consequentemente sua opacidade.

O amido produz filmes mais opacos conforme observado por Fakhoury et al.[10], quando comparou diferentes proporções de amido e gelatina em filmes obtidos por casting e plastificados com glicerol, em que com o aumento do percentual de amido em relação a gelatina foi observado um aumento da opacidade, de aproximadamente 7 a 24% com a relação amido/gelatina passando de 1/4 para 4/1 respectivamente.

O mesmo comportamento foi observado por Garcia et al.[11] em filmes de amido e quitosana, onde o amido promoveu um aumento da opacidade dos materiais.

Rocha et al.[12] produziram filmes de amido de mandioca e proteína de soja plastificados com glicerol pela técnica de casting e observaram que com o aumento da concentração de proteína de soja houve aumento da opacidade aparente e do ΔE* (variando de 8,14 a 17,62) dos filmes, e conforme os autores essa diferença foi devido à coloração amarelada característica da proteína de soja. Almeida et al.[13] produziram filmes de amido de batata e celulose microbiana plastificados com glicerol pela técnica de casting e observaram uma maior opacidade nos filmes em função do aumento da concentração de celulose microbiana nos filmes. Ambos os autores supracitados observaram que o amido não foi um fator importante para a opacidade e o ΔE* dos filmes, comportamento contrário ao observado no presente trabalho e essa diferença pode ser atribuída ao processo de produção dos materiais. A técnica de casting tende a produzir materiais mais translúcidos de que aqueles obtidos por extrusão, onde podem ocorrer reações de degradação do amido ou do PVA devido às altas temperaturas e taxas de cisalhamento, tornando os materiais com coloração mais escura e mais opaca de que filmes produzidos por casting.

A espessura também é um fator determinante para o aumento da opacidade, Mali et al.[14], observaram um aumento da opacidade em filmes de amido de cará plastificados com glicerol com o aumento da espessura dos mesmos, sendo os materiais obtidos no presente estudo apresentaram espessura média de 858 µm, promovendo assim um aumento esperado na opacidade e diferença de cor.

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4. Conclusões Os laminados biodegradáveis a base de amido e PVA obtidos por extrusão se mostraram levemente opacos, sendo o PVA e a interação amido/PVA os principais responsáveis pelo aumento da opacidade dos laminados. A diferença de cor é devida principalmente ao efeito do amido, que promoveu materiais de coloração mais amarelada.

5. Agradecimentos Os autores agradecem a Universidade Tecnológica Federal do Paraná, ao CNPq, CAPES e Fundação Araucária pelas bolsas de estudo e apoio financeiro.

6. Referências 1. Liu, H., Xie, F., Yu, L., Chen, L., & Li, L. (2009). Thermal processing of starch-based Polymers. Progress in Polymer Science, 34(12), 1348-1368. http://dx.doi.org/10.1016/j. progpolymsci.2009.07.001.

2. Olivato, J. B., Müller, C. M. O., Yamashita, F., Grossmann, M. V. E., & Nobrega, M. M. (2013). Study of the compatibilizer effect in the properties of starch/polyester blends. Polímeros: Ciência e Tecnologia, 23(3), 346-351. http://dx.doi.org/10.4322/ polimeros.2013.014.

3. Brandelero, R. P. H., Yamashita, F., & Grossmann, M. V. E. (2010). Effect of the method of production of the blends on mechanical and structural properties of biodegradable starch films produced by blown extrusion. Carbohydrate Polymers, 82(4), 1102-1109. http://dx.doi.org/10.1016/j.carbpol.2010.06.034.

4. Maria, T. M. C., Carvalho, R. A., Sobral, P. J. A., Habitante, A. M. B. Q., & Solorza-Faria, J. (2008). The effect of the degree of hydrolysis of the PVA and the plasticizer concentration on the color, opacity, and thermal and mechanical properties of films based on PVA and gelatin blends. Journal of Food Engineering, 87(2), 191-199. http://dx.doi.org/10.1016/j. jfoodeng.2007.11.026.

5. Pelissari, F. M., Yamashita, F., Garcia, M. A., Martino, M. N., Zaritzky, N. E., & Grossmann, M. V. E. (2012). Constrained mixture design applied to the development of cassava starch – chitosan blown films. Journal of Food Engineering, 108(2), 262-267. http://dx.doi.org/10.1016/j.jfoodeng.2011.09.004.

6. Jang, J., & Lee, D. (2003). Plasticizer effect on the melting and crystallization behavior of polyvinyl alcohol. Polymer, 44(26), 8139-8156. http://dx.doi.org/10.1016/j.polymer.2003.10.015.

7. Sobral, P. J. A. (2000). Influência da espessura de biofilmes feitos à base de proteínas miofibrilares sobre suas propriedades funcionais. Pesquisa Agropecuaria Brasileira, 35(6), 1251-1259. http://dx.doi.org/10.1590/S0100-204X2000000600022.

8. Moraes, I., Silva, G. G. D. I., Carvalho, R. A., Habitante, A. N. Q. B., Bergo, P. V. A., & Sobral, P. V. A. (2008). Influência do grau de hidrólise do poli(vinil álcool) nas propriedades físicas de filmes à base de blendas de gelatina e poli(vinil álcool) plastificados com glicerol. Ciência e Tecnologia de Alimentos, 28(3), 738-745. http://dx.doi.org/10.1590/S010120612008000300034.

9. Silva, G. G. D. I., Sobral, P. J. A., Carvalho, R. A., Bergo, P. V. A., Mendieta-Taboada, O., & Habitante, A. M. Q. B. (2008). Biodegradable films based on blends of gelatin and poly (vinyl alcohol): effect of PVA type or concentration on some physical properties of films. Journal of Polymers and the Environment, 16(4), 276-285. http://dx.doi.org/10.1007/ s10924-008-0112-9.

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Laminados biodegradáveis de blendas de amido de mandioca e poli(vinil álcool): efeito da formulação sobre a cor e opacidade 10. Fakhoury, F. M., Martelli, S. M., Bertan, L. C., Yamashita, F., Mei, L. H. I., & Queiroz, F. P. C. (2012). Edible films made from blends of manioc starch and gelatin – Influence of different types of plasticizer and different levels of macromolecules on their properties. LWT – Food Science and Technology, 49(1), 149-154. http://dx.doi.org/10.1016/j.lwt.2012.04.017.

11. Garcia, M. A., Pinotti, A., & Zaritzky, N. (2006). Physicochemical, water vapor barrier and mechanical properties of corn starch and chitosan composite films. Starch, 58(9), 453-463. http:// dx.doi.org/10.1002/star.200500484.

12. Rocha, G. O., Farias, M. G., Carvalho, C. W. P., Ascheri, J. L. R., & Galdeano, M. C. (2014). Filmes compostos biodegradáveis a base de amido de mandioca e proteína de soja. Polímeros: Ciência e Tecnologia, 24(5), 587-595. http:// dx.doi.org/10.1590/0104-1428.1355.

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13. Almeida, D. M., Woiciechowski, A. L., Wosiacki, G., Prestes, R. A., & Pinheiro, L. A. (2013). Propriedades físicas, químicas e de barreira em filmes formados por blenda de celulose bacteriana e fécula de batata. Polímeros: Ciência e Tecnologia, 23(4), 538-546. http://dx.doi.org/10.4322/polimeros.2013.038.

14. Mali, S., Grossmann, M. V. E., Garcia, M. A., Martino, M. N., & Zaritzky, N. E. (2004). Barrier, mechanical and optical properties of plasticized yam starch films. Carbohydrate Polymers, 56(2), 129-135. http://dx.doi.org/10.1016/j.carbpol.2004.01.004.

Enviado: Jun. 25, 2014 Revisado: Jan. 12, 2015 Aceito: Fev. 04, 2015

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