Polímeros Ciência e Tecnologia, vol.25, n.6, 2015 by Polímeros: Ciência e Tecnologia (Polimeros)

A SABIC é líder mundial na produção de resinas termoplásticas de alto desempenho e conta hoje com um dos mais amplos portfólios de produtos do mundo. Com duas fábricas na América do Sul (Brasil e Argentina), a SABIC possui uma equipe técnica experiente, pronta para auxiliar seus clientes no desenvolvimento de novos projetos com resinas termoplásticas.

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http://dx.doi.org/0.1590/0104-1428.2506

Editorial

Já há algum tempo venho atuando como membro do Comitê Editorial da Polímeros, o que me tem dado grande

Caros leitores da revista Polímeros satisfação. Recentemente fui procurado pelo Prof. Luiz A. Pessan, Presidente da ABPol que falando também em nome do Prof. Marco-Aurelio De Paoli (UNICAMP/IQ), Presidente do Conselho Editorial desta revista me convidavam para assumir a posição de Presidente de Comitê Editorial. Reconhecedor do excelente trabalho que vinha sendo conduzido tarefa. Espero estar à altura dos colegas que ocuparam este posto antes de mim, e não dar tanto trabalho ao próximo.

pela Profa. Elisabete Frollini (USP/IQSC) durante seu mandato hesitei, mas após várias reuniões aceitei assumir tão nobre Meu aceite se embasou na confiança que tenho nos Membros do Comitê Editorial os professores Adhemar C. Ruvolo Filho (UFSCar/DQ), Bluma G. Soares (UFRJ/IMA), César Liberato Petzhold (UFRGS/IQ), Glaura Goulart Silva (UFMG/DQ), José António C. Gomes Covas (UMinho/Portugal), José Carlos C. S. Pinto (UFRJ/COPPE) e Regina do Prof. João B.P. Soares da University of Alberta, Canadá, que entusiasticamente aceitou nosso convite. Confiança

Célia R. Nunes (UFRJ/IMA). Para reforçar ainda mais “este time” brevemente também contaremos com a participação também nos funcionários da secretaria da ABPol Sr. Marcelo P. Gomes e especialmente Charles F. de Souza, Assistente Editorial desta revista. Assim é com muito prazer que assumirei esta função a partir de Julho de 2016, compromissos profissionais me C. Rúvolo Filho, que exercerá a função de Editor Interino. Caro Adhemar, obrigado pela cooperação, sei que você não

impedem que eu assuma imediatamente. Durante este período de transição contarei com a valiosa ajuda do Prof. Adhemar mede esforços em prol de nossa querida “revista”. Voltarei a ocupar este espaço a partir das edições do segundo semestre de 2016 quanto então assumirei em definitivo a função de Editor Presidente do Comitê Editorial. Meus agradecimentos aos Membros do Comitê Editorial e em especial Á você caro leitor desejo um Feliz 2016, com muita saúde e disposição para cooperativamente contribuirmos com

ao amigo Adhemar. a produtividade científica que o Brasil tanto necessita.

Sebastião V. Canevarolo Jr. Editor

Boa leitura a todos

Polímeros, 25(6), 2015

P o l í m e r o s - N º 6 - V o l u m e X X V - N o v / D e z - 2 0 1 5 - ISS N 0 1 0 4 - 1 4 2 8

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

Polímeros P r e s i d en t e

Conselho Editorial

Marco-Aurelio De Paoli (UNICAMP/IQ)

Membros

Conselho Editorial

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) João B. P. Soares (UAlberta/DCME) José Alexandrino de Sousa (UFSCar/DEMa) José António C. Gomes Covas (UMinho/IPC) José Carlos C. S. Pinto (UFRJ/COPPE) Júlio Harada (Harada Hajime Machado Consutoria Ltda) Laura H. de Carvalho (UFCG/DEMa) Luiz Antonio Pessan (UFSCar/DEMa) Luiz Henrique C. Mattoso (EMBRAPA) 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)

Comitê Editorial Sebastião V. Canevarolo Jr. – Editor

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Comitê Editorial

Adhemar C. Ruvolo Filho Bluma G. Soares César Liberato Petzhold Glaura Goulart Silva João B. P. Soares José António C. Gomes Covas José Carlos C. S. Pinto Regina Célia R. Nunes

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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

Bimestral v. 25, nº 6 (nov./dez. 2015) ISSN 0104-1428

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

1. Polímeros. l. Associação Brasileira de Polímeros. E2

Polímeros, 25(6), 2015

Polímeros Seção Editorial

Editorial................................................................................................................................................................................................E1 Informes & Notícias ............................................................................................................................................................................E4 Calendário de Eventos ........................................................................................................................................................................E5 Associados...........................................................................................................................................................................................E6

S e ç ã o T é cn i c a Photodegradation of a polypropylene filled with lanthanide complexes Valérie Massardier and Molka Louizi.............................................................................................................................................................. 515 Robson Fleming Ribeiro, Luiz Claudio Pardini, Nilton Pereira Alves and Carlos Alberto Rios Brito Júnior................................................ 523

Thermal Stabilization study of polyacrylonitrile fiber obtained by extrusion The influence of long chain branches of LLDPE on processability and physical properties Paula Cristina Dartora, Ruth Marlene Campomanes Santana and Ana Cristina Fontes Moreira................................................................. 531

Surface treated fly ash filled modified epoxy composites Uma Dharmalingam, Meenakshi Dhanasekaran, Kothandaraman Balasubramanian and Ravichandran Kandasamy................................ 540 Manuel Rapado and Carlos Peniche............................................................................................................................................................... 547

Synthesis and characterization of pH and temperature responsive poly(2-hydroxyethyl methacrylate-co-acrylamide) hydrogels Molecular weight and tacticity effect on morphological and mechanical properties of Ziegler–Natta catalyzed isotactic polypropylenes Ismael Amer, Albert van Reenen and Touhami Mokrani.................................................................................................................................. 556

Design of conformal cooling for plastic injection moulding by heat transfer simulation Sabrina Marques, Adriano Fagali de Souza, Jackson Miranda and Ihar Yadroitsau...................................................................................... 564 Marcelo Aparecido Chinelatto, José Augusto Marcondes Agnelli and Sebastião Vicente Canevarolo........................................................... 575

Biobased additive plasticizing Polylactic acid (PLA)

Synthesis and photostabilizing performance of a polymeric HALS based on 1,2,2,6,6-pentamethylpiperidine and vinyl acetate

Mounira Maiza, Mohamed Tahar Benaniba, Guilhem Quintard and Valerie Massardier-Nageotte.............................................................. 581

Characterization of clay filled poly (butylene terephthalate) nanocomposites prepared by solution blending Khalid Saeed and Inayatullah Khan................................................................................................................................................................ 591 Layla Talita de Oliveira Alves, Cynthia D’Avila Carvalho Erbetta, Christian Fernandes, Maria Elisa Scarpelli Ribeiro e Silva, Roberto Fernando Souza Freitas e Ricardo Geraldo Sousa............................................................................................................................ 596

Síntese e caracterização de MIP com fenilalanina visando sua aplicação na técnica de SPE

Incorporação de fonte de nitrogênio em partículas de PVA e alginato de sódio e estudo da influência de ciclos de congelamento/ descongelamento na caracterização do produto Sinara Queli Welter Nardi, Sirlei Dias Teixeira e Cristiane Regina Budziak Parabocz.................................................................................. 606

Capa: Cross section SEM images of fracture morphologies of extruded PAN co-VA fibers. Micrografias do MIP-1 antes da extração da Phe (a) e após a extração (b). SEM micrographs of isotactic polypropylene fractions: (A) P5(120) (Mw = 207823 g/mol); (B) P4(120) (Mw = 195693 g/mol) and (C) P9(110) (Mw =110387 g/mol) (3000x magnification). SEM micrographs of isotactic polypropylene polymers: (A) P5 (mmmm = 96%); (B) P4 (mmmm = 94%) and (C) P14 (mmmm = 86%) (3000x magnification). Elaboração artística Editora Cubo.

Polímeros, 25(6), 2015

I N F O R M E S

Biobased polymer market, “slow-growth in 2015” but triples by 2020: European Bioplastics report

In Europe, the European Bioplastics trade group projected that worldwide production capacity of bio‑based polymer will triple from 5.7 million tonnes in 2014 to nearly 17 million tonnes in 2020. The data show a 10% growth rate from 2012 to 2013 and even 11% from 2013 to 2014. However, growth rate is expected to decrease in 2015. Consequence of the low oil price? The third edition of the well-known 500 pagemarket study and trend reports on “Bio-based Building Blocks and Polymers in the World – Capacities, Production and Applications: Status Quo and Trends Towards 2020” was published last week, and includes consistent data from the year 2012 to the latest data of 2014 and the recently published data from European Bioplastics, the association representing the interests of Europe’s bioplastics industry. Bio-based drop-in PET and the new polymer PHA show the fastest rates of market growth. Europe looses considerable shares in total production to Asia. The bio-based polymer turnover was about €11 billion worldwide in 2014 compared to €10 billion in 2013. Source: Biofuels Digest

Polymer holds hope for jute sector

N O T Í C I A S

Scientists have developed a jute-based polymer that could be used to manufacture under-the-hood automobile parts - like air intake manifold, radiator end-caps, fan and shroud - or housing construction or even microwavable cooking containers. It could revolutionize the car industry by making parts cheaper and inject a massive boost to the dying jute industry. The breakthrough has been achieved at Bengaluru‑based Steer, which works in the field of pharmaceuticals, plastics, food and nutraceuticals, biomaterials and bio-refining. Its scientists say the jute-filled polypropylene (PP) compound they have developed by incorporating up to 50% jute (by weight) is a strong, heat-resistant, flexible, light, economical and eco-friendly reinforcing agent for plastics. “The material

can absorb up to 150°C, take any form and shape, and any colour. There is huge scope in automobile industry as it can be moulded into a car’s bumpers or dashboards,” said Steer director Shyam Sunder Pani. West Bengal jute commissioner Subrata Gupta termed the development significant and felt it could revive the beleaguered jute industry, which is gasping under the choke-hold of plastic bags. Bengal, the largest producer of jute, witnesses periodic shutdowns of mills due to falls in jute demand. Of the 94 composite jute mills in the country, 74 are in Bengal. India produces 1.5 million tonne of jute but it has been languishing at the low end of the value chain of the agriculture produce pyramid while products like sugarcane, coconut, coffee and tea which have become highly commercialized commodities, delivering huge monetary benefits to producers and value adders alike. According to Steer founder and managing director Babu Padmanabhan, the multiple benefits of jute polymers will fuel demand and jobs, thereby reviving the industry. “Jute has for decades been on the decline and considered a sunset industry. But with jute polymers, there is a potential for a new sunrise industry to emerge, creating thousands of jobs in jute rich resource states like Bengal, Bihar and Odisha.” According to government estimates, the jute industry provides direct employment to nearly 0.37 million workers and supports the livelihood of around 4 million farm families; in addition to the large number of people engaged in the trade of the natural fiber. Jute polymers, scientists say, will not only help reduce product cost, density and carbon footprint, it will also enhance product performance. It will transform the jute industry where nearly 75% jute goods are currently used as packaging materials, burlap, gunny cloth and sacks. Hemant Bangur, chairman of Fort Gloster Jute Mills, arguably the most modern jute mill in India, termed it a very interesting development and good for the survival of the industry. “A lot of car makers in Europe are already using jute polymers. For it to start here, we have to work in tandem with automobile companies. Besides, we need regulatory protection. If it is implemented, many auto companies will be keen to use this material,” he said. Source: Bennett, Coleman & Co. Ltd

Polímeros, 25(6), 2015

May ExpoPlast Perú 2016 Date: 3–6 May 2016 Local: Lima- Peru Website: http://www.expoplastperu.com/ Guangzhou International Wood-Plastic Composites Fair 2016 Date: 13–16 May 2016 Local: Guangzhou - China Website: http://www.musuz.com/ 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

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/

August Interplast 2016 Date: 16–19 August 2016 Local: Joinville - SC Website: www.messebrasil.com.br

September Polycondensation 2016 Date: 11–15 September 2016 Local: Moscow / St Petersburg - Russian Website: http://www.polycondensation2016.ac.ru/index.php/en/

June

Polyolefin Additives – 2016 Date: 13–15 September 2016 Local: Vienna - Austria Website: http://www.amiplastics.com/events/event?Code=C743

Polymer Compounding for Innovations in Plastics Industry Date: 7–9 June 2016 Local: Newark - USA Website: http://www.compoundingconference.com/

PLASTEC Minneapolis Date: 21–22 September 2016 Local: Minnesota - USA Website: http://plastecminn.plasticstoday.com/

COMPLAST Kenya Plast 2016 Date: 8–10 June 2016 Local: Nairobi - Kenia Website: http://www.kenyaplast.in/

Organic Semiconductors Date: 22–25 September 2016 Local: Dubrovnik - Croatia Website: http://www.zingconferences.com/conferences/organicsemiconductors/

Argenplás 2016 Date: 13–16 June 2016 Local: Buenos Aires - Argentina Website: http://www.argenplas.com.ar/ Oil & Gas Polymer Engineering Texas 2016 Date: 14–15 June 2016 Local: Texas - USA Website: http://www.amiplastics.com/events/event?Code=C734 Plastics Design and Moulding 2016 Date: 14–15 June 2016 Local: Telford - UK Website: http://www.pdmevent.com/ PLASTEC East Date: 14–16 June 2016 Local: New York - USA Website: http://plastec-east.plasticstoday.com/ Propak Asia 2016 Date: 15–18 June 2016 Local: Bangcoc - Thailand Website: http://www.propakasia.com/ Polymers in Cables – 2016 Date: 21–22 June 2016 Local: Pennsylvania - USA Website: http://www.amiplastics.com/events/event?Code=C732

Colombiaplast 2016 Date: 26–30 September 2016 Local: Bogotá - Colombia Website: http://www.colombiaplast.com/ Conductive Plastics - 2016 Date: 26–30 September 2016 Local: Pennsylvania - USA Website: http://www.amiplastics.com/events/event?Code=C742

October Polymeric Implants & Catheters in Medical Devices Date: 4–6 October 2016 Local: Las Vegas - USA Website: http://www.mediplastconference.com/ China International Exhibition on Plastics and Rubber Injection Moulding Industry (CIM) 2016 Date: 13–15 October 2016 Local: Tianjin - China Website: http://www.cimexpo.cn/

November

July

Polymer Foam – 2016 Date: 8–10 November 2016 Local: Cologne - Germany Website http://www.amiplastics.com/events/event?Code=C752

COMPLAST Plastics Myanmar 2016 Date: 08–10 July 2016 Local: Yangon - Myanmar Website: http://www.plastics-myanmar.in/

Expoplast 2016 Date: November 30 - December 1, 2016 Local: Québec - Canada Website: http://expoplast.plasticstoday.com/

Polímeros, 25(6), 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

Polímeros, 25(6), 2015

Associados da ABPol 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(6), 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/10.1590/0104-1428.2006

Photodegradation of a polypropylene filled with lanthanide complexes Valérie Massardier1*and Molka Louizi1 UMR 5223 Ingénierie des Matériaux Polymères, Centre National de la Recherche Scientifique – CNRS, Institut National des Sciences Appliquées de Lyon – INSA Lyon, Villeurbanne, France

*valerie.massardier@insa-lyon.fr

Abstract This research aims at studying the photodegradation of a polypropylene formulation filled with lanthanide complexes. These complexes can be used as tracers for the identification of polymer materials in order to facilitate an high speed automatic sorting of plastic wastes for an economically efficient recycling. By paying attention to the evolution of carbonyl absorption bands in FTIR spectra, it is observed that the addition of lanthanide complexes into our formulation improves UV stability of polypropylene by reducing the extent of photo-degradation. Furthermore, TG analyses show that the traced blends can maintain better thermal properties, after irradiation. A significant increase of the crystallinity degree and a decrease of the melting temperature are more pronounced for the unfilled UV–irradiated PP. This might result from chemi-crystallization that can occur when chain entanglements are broken as a result of chain scissions. From SEM analyses, it is observed that the severity of surface cracks induced by photo degradation is reduced for filled PP. The mechanical tests are in agreement with this result and show a fundamental change in the behavior of the as-exposed blends from a ductile to a brittle material. Keywords: polypropylene, lanthanide complexes, UV irradiation, photodegradation.

1. Introduction The commercial importance of polymers has lead to very large applications in the form of composites/nanocomposites, polymer blends[1-4], such as automotive, aerospace, packaging, etc. Whatever the application and especially outdoor ones, there is often a natural concern regarding the durability of polymer materials. Commonly known, the durability of any material depends on several factors such as the environment (especially sunlight intensity, temperature, moisture, etc.), the exposure time, the type of polymer (presence of thermal and UV stabilizers, etc.), water absorption, etc[5]. Among these factors, ultra violet (UV) irradiation is a frequently encountered factor that can induce photo-degradation of polymers. Polypropylene (PP) is one of the most used polyolefins, with a significant part of its applications corresponding to outdoor environments. For this reason, a lot of studies are focused on the understanding of the mechanisms of UV‑induced degradation of PP[6-9] and polyolefins[10,11]. Photolytic degradation and photo-oxidation are the most important phenomenon observed when irradiating PP within the active range [310-350 nm][5,12]. These reactions preferentially occur in the amorphous fraction because of its high permeability to oxygen[13]. Photolysis primarily involves the adsorption of wavelengths above 290 nm in the UV regions of the solar spectrum by chromophores (or impurities) which in turn release sufficient energy to cause bond scissions resulting in the formation of radicals that can either combine to form more chromophores or initiate photo-oxidation[14]. It is worth noting that these mechanisms occur in the UV degradation of PP and lead to main changes in their chemical, physical and mechanical properties. It is generally known that the photodegradation

Polímeros, 25(6), 515-522, 2015

induces a change in the melting and cristallinity behaviors of PP. The melting temperature decrease is ascribed to i) oxidative reactions on crystal surface that increase the surface free energy of the crystals and ii) new crystals formed from polymer released by degradation that may have a lower melting temperature than the pre-existing crystals because of the defects content (crosslinks, carbonyl groups, etc)[15]. As regards the cristallinity, the freed segments, formed by molecular chain scissions, form a new crystalline structure in the amorphous zone, especially at high exposure temperature (about 65°C) and thus provoke the growth of pre-existing crystals of PP[16]. This increase in crystallinity is a form of secondary crystallization often known as chemi-crystallization[17]. It should be noted that the increase in crystallinity is limited by the presence of chemical irregularities such as carbonyls and hydroperoxides that form progressively under photooxidation. Thus, there are two opposite effects, one that promotes greater crystallinity and the other inhibiting crystallization. The shorter chains produced by scission events will crystallize more rapidly whereas crosslinks and molecular defects (carbonyls, etc.) will not be able to crystallize and will be rejected from the newly formed crystals. In addition, the phenomena occurring during the PP photodegradation (chain scissions and in turn shrinkage) lead to the development of tensile residual stresses near the surface[6]. A severe deterioration of engineering properties may therefore occur. The most important practical consequence of chemi-crystallization is the formation of cracks caused by contraction of the surface layers[18]. The presence of surface cracks is indeed one of the main reasons for the embrittlement of ductile semi‑crystalline polymers such as PP, causing serious deterioration in the

515

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Massardier, V., & Louizi, M. mechanical properties of the polymer after short-term exposures[16,19]. In order to decrease the UV degradation and extend the lifetime of PP polymer, there is a growing interest in adding UV screeners to bulk polymers during processing. The photolysis of PP is induced by adsorption of UV light causing the formation of free radicals. Those radicals can react with oxygen to yield hydroperoxides and then carbonyl groups[12,14]. There are also several studies focusing on the thermal stability and photo-degradation mechanisms for the obtained nanocomposites. Carbon black (CB) has been considered as the most effective organic UV screener as it absorbs all incident lights with its phenolic and quinoid groups functioning as anti-oxidants[20]. As regards the inorganic compounds, the zinc oxide proved to have a relatively high adsorption band starting at 385 nm and extending to the far-UV[21]. It is shown that zinc oxide, as a UV-screening additive, does not degrade when absorbing light and in many cases may improve mechanical, optical and electronic properties[12]. In terms of UV damage, it is also found that layered silica filled polymer composites exhibit remarkable improvement of mechanical, thermal properties when compared with pure polymer due to interactions of fillers with the polymer matrix at the nano-level scale[22]. Zhao and Li[23] show that the extent of photodegradation is significantly reduced with the addition of ZnO nanoparticles in PP matrix. This is ascribed to the superior UV light screening effects offered by the ZnO nanoparticles. Despite several researches focused on the study of the photo-degradation of polymers filled with nanosilica, zinc oxide, carbon black, MMT, etc only few studies deal with polymers containing rare earth oxides[24-26]. Recently, Bezati et al.[27] have shown that the addition of 1 wt% of cerium oxide nanoparticles improves the photo‑degradation resistance of PP matrix to UV exposure due to the light screening effects offered by these particles. As outlined above, the oxidation of PP chains can be reduced by adding specific fillers (nanosilica, zinc oxide, carbon black, MMT, rare earth oxide, etc) acting as UV filters and hence protects the polymer from UV intensity. Therefore, in the present article, we are interested in the photodegradation of polypropylene filled with lanthanide complexes that can be used as tracers for the identification of polymer materials (especially black ones), in order to facilitate high speed automatic sorting of plastic wastes for an economically efficient recycling. The feasibility of the detection of these lanthanide complexes dispersed in white and black PP matrices was successfully achieved in previous studies[28] through the use of UV-ray fluorescence spectrometry. Furthermore, the elaboration of traced blends

via high shear process[29,30] has proved to be a major key for preparation of well dispersed tracer particles within polypropylene matrix especially when processing at a screw speed, N, equal to 800 rpm. The present study was conducted to gain a good knowledge of the subsequent changes occurring in the properties of traced black polypropylene with lanthanide complexes under UV irradiation. In our case, for accelerating ageing phenomenon, the surface of the traced blends was UV-irradiated at 60°C for 3 months (2160 hours). The effectiveness of lanthanide complexes particles on combating UV irradiation damage for PP is evaluated through various analyses to characterize evolutions in chemical structures (FTIR), in morphologies (MEB), in thermal properties (TGA, DSC), as well as in mechanical properties (tensile and impact tests).

2. Materials and Methods 2.1 Materials The polymer investigated in this study is a commercial grade, BMU 133, used in the automotive as well as in the electrical & electronic fields. BMU 133 is a black polypropylene copolymer provided by Exxon Mobil Chemicals. It contains 15 wt% of elastomer and 10 wt% of carbon black (CB) shows a Melt Flow Index MFI = 15 g/10 min under 2.16 kg at 230°C and a specific gravity = 0.970 g/cm3. Two fluorescent tracers (T2 and T3) were tested and belong to the inorganic family of the lanthanide complexes. T2 is a doped aluminum and barium oxide, T3 is a doped vanadium trioxide. These tracers do not chemically react with the materials and were provided by the start up Tracing Technologies. They are thermally stable at high temperatures, compatible with the REACH regulations. Their essential feature is that when they are excited with a UV source between 300 and 400 nm, they fluoresce in the visible light spectrum.

2.2 Elaboration of traced blends All blends filled with 0.1 wt% of fluorescent tracers (T2 or T3) were prepared with a co-rotating twin screw‑extruder (TSE) Leistritz ZSE 18 HP. The screw profile, with a screw diameter of 18 mm and a L/D ratio of 60, used for all the experiments is illustrated in Figure 1. The originality of this extruder lies on reaching high shear rates on extruded blends by adjusting the screw rotation speed. In addition, in the first part of this study, this technique has proved to be a key for the preparation of well-dispersed tracer

Figure 1. Screw configuration of the co-rotating high shear extruder (Leistritz ZSE 18 HP). 516

Polímeros , 25(6), 515-522, 2015

Photodegradation of a polypropylene filled with lanthanide complexes particles within polypropylene matrix. The screw speed, the mean residence time and the feed rate (Q) were 800 rpm, 55 s, and 3 kg/h respectively. The chosen rotation speed, 800 rpm, corresponds to a shear rate of 500 s–1 calculated thanks to Ludovic software[31]. The extrusion temperature was fixed at 200°C but the experimental melt temperature at the die exit was about 235°C, measured by introducing a thermocouple into the bulk extrudate. All traced blends were prepared under identical mixing and moulding conditions. At the die exit of the extruder, the extrudates are pelletized and then moulded at 200°C through a Battenfeld 350 PLUS injection moulding machine. Standard tensile and rectangular bars are produced for mechanical analyses.

2.3 UV irradiation As regards the photodegradation of the traced blends, a UV irradiation treatment was accomplished for the injected samples using a QUV accelerated weathering machine (Q-panel lab products, UVA 340 nm), with a light intensity of 0.68 W/m2. QUV accelerated weathering was used to simulate long-term exposure. The source of UV irradiation was fluorescent tubes UVA-340, with an output responsibly close to the solar radiation in the UV range (Labomat). The surface of samples was UV irradiated at a controlled temperature of 60°C for an exposure time of 3 months (2160 hours). The traced blends containing 0.1 wt% of tracers (T) and UV irradiated will be designed as UV‑ (for example: UV‑BMU 133-T3-0.1) and the non-irradiated will be referred as polymer type-T-0.1 (for instance: BMU 133-T3-0.1).

2.4 Infrared spectrometry UV irradiated traced blends were characterized by Fourier Transform Infrared spectrometry (FTIR). Samples of about 0.4 mm thick were cut from the surface of the plates and observed. The FTIR-ATR measurements were obtained in absorption mode by using 16 scans at 2 cm–1. PP molecular degradation is characterized by calculation of a carbonyl index (CI), which was calculated by the following equation: CI =

Ac (1) AR

Where Ac represents the area of the carbonyl absorption band being in the range [1700-1800 cm–1] and AR is the area of the reference band in the range [2700-2750 cm–1][23]. The reference peak corresponds to the CH3 stretching and CH bending[16].

2.5 Thermal properties The thermal properties of blends before and after UV irradiation were carried out by Differential Scanning Calorimetry by using DSC Q10 of TA instruments. To perform these tests, samples were cut into pellets and placed in aluminum pans. A scan was performed from 10 to 200°C, further maintained for 2 min at 200°C to erase the thermal history of the blends and then cooled down from 200°C to 10°C. The heating and cooling rates were fixed at 10°C/min. The crystallization degree, XC, is calculated by considering a melting enthalpy of 209 J/g for a 100% crystalline polypropylene. Polímeros, 25(6), 515-522, 2015

To characterize the thermal stability of samples before and after UV irradiation, thermogravimetric analyses with TA Q500 apparatus were carried out. The measurements were conducted under an argon flow rate at heating rates of 20°C/min. The scanning temperature was in the range [25-600°C]. The temperature of maximum decomposition rate was determined for all the blends. It is worth noting that after UV treatment, prior to DSC and TGA analyses, thin slices of samples, about 0.4 mm thick, were cut from surfaces of the molded plates for the measurements.

2.6 Morphology observation The dispersion of tracers in the selected polymers was investigated by Scanning Electron Microscopy (SEM) with a Hitashi S800 model at an accelerating voltage of 30 kV. Prior to observations, samples were cryo-fractured in liquid nitrogen to avoid any plastic deformation. SEM was also used to observe the surface cracking of samples after UV irradiation.

2.7 Mechanical tests Tensile tests were carried out by means of an Instron machine MTS 2/M tester, at a crosshead speed of 30 mm/min at room temperature. Impact tests were performed by means of a Zwick D7900 Type 5102-100/00 instrument in compliance with standard ISO 179 on notched specimens conditioned at –22°C for 48h. These tests were conducted in order to obtain the Young’s Modulus, elongation at break, tensile and impact strengths before and after irradiation treatment. All the reported values are the averages of ten experimental results to check the reproducibility.

3. Results and Discussions 3.1 Evolution of chemical structures As shown previously, the photodegradation process of materials corresponds to radical reactions in chains that can be initiated by the presence of impurities and high energy photon collision[23]. By using FTIR spectroscopy, it is easy to identify the degradation products generated by macroradical oxidation. In previous researches[32,33], it was shown that the photo-degradation of PP causes mainly the formation of hydroperoxides and carbonyls easily observed in the wavenumber ranges [3200-3600 cm–1] and [1700-1800 cm–1] respectively. In this study, evolutions of the FTIR spectra of UV-BMU133, UV-BMU133-T2 or T3 were examined in the wavenumber range [1700-1800 cm–1] (Figure 2) and the carbonyl index of these samples after UV-treatment was also calculated by using Equation 1 and displayed in Figure 3. The spectrum of unfilled BMU133 before irradiation was also carried out for comparison. As expected, the UV-BMU133 blend presents an intense peak in the carbonyl region [1700-1800 cm–1] after UV irradiation. For the irradiated samples containing T3 or T2 tracer, the carbonyl peak intensity is decreased as well as the carbonyl index compared to UV-BMU133 blends (Figures 2 and 3). This can be ascribed to the effect of tracer particles which 517

Massardier, V., & Louizi, M.

Figure 2. FTIR spectra of UV irradiated BMU133 and its traced blends with T2 or T3.

3.2 Effect of UV irradiation on the thermal properties of traced blends

Figure 3. Carbonyl index of UV-BMU133, UV-BMU133-T2-0.1 or T3-0.1.

contributes to the stabilization of PP chains and delay the photodegradation process by acting as screens. To play this role, tracer particles must absorb a part of UV irradiation and thus contribute to the decrease of UV intensity that can promote the oxidation of the PP chains. Furthermore, the mechanism of photo degradation of UV‑BMU133 without or with T2 or T3 tracer after irradiation is identical and no additional peak appears when adding tracers. These findings are in good agreement with those observed by Bezati et al. [27] for PP samples containing rare earth particles exposed to UV irradiation. Similar stabilization effect is observed for metal oxides by several authors[25,26]. 518

To have an insight on the effect of UV irradiation treatment on thermal properties of unfilled BMU133 and its traced blends, DSC and TGA analyses were carried out. The results in terms of melting/crystallization temperatures and enthalpies as well as the temperature at maximum weight loss (Tmax) of these samples are summarized in Table 1. On a one hand, it is well observed that the UV irradiation treatment has a considerable effect on the crystallization and melting behaviors of the treated samples (Table 1). The melting and crystallization temperatures of UV irradiated BMU133 shift to lower temperatures from 168.2 to 151.5°C for Tm and from 124.0 to 112.8°C for Tc. Furthermore, a reduction in Tm is shown for all the UV irradiated traced blends due to the UV irradiation. Such an effect is expected because the photo-degradation causes chain scissions resulting in the formation of more freed segments[16,23]. This contributes to a drop off of molecular weight of materials and consequently leads to a decrease of the melting temperature[27] . The decrease in melting temperature can also be due to oxidative reactions on the crystal surface that increase the surface free energy of the crystals[15,34]. Similar trends were observed by Bezati et al.[27] and Rabello and White[16] in previous studies. From Table 1, it can also be noted that the cristallinity (Xc) of BMU133 is increased upon UV irradiation. The change in crystallization behaviour can be explained by the chemi-crystallization occurring when chain entanglements are broken as a result of chain scissions. Then, the freed segments produced by molecular chain scissions can induce the formation of new crystalline structures in the amorphous zone, especially at high exposure temperature (about 65°C) and thus provoke the growth of pre-existing crystals of PP[16]. Furthermore, it Polímeros , 25(6), 515-522, 2015

Photodegradation of a polypropylene filled with lanthanide complexes is worth noting that the presence of chemical irregularities such as carbonyls and hydroperoxides that form progressively under photo-oxidation can also contribute to a disruption in crystallinity as shown by Zhao and Li[23]. Regarding the crystallisation degrees (Xc ) of traced blends, they are not significantly affected in presence of lanthanide complexes. This can be due to the stabilization effect of tracer particles that blocks the chemi-crystallization of PP[16]. On the other hand, it is shown that UV-irradiation impacts the thermal stability of BMU133 as well as its traced blends with T2 and T3 tracers. As reported in Table 1, the thermal decomposition of UV-BMU133 occurs at lower temperatures after irradiation compared to those being non-irradiated. This is expected because when polymer formulations are irradiated with ultraviolet radiations, they can degrade due to absorption of light energy by chemical groups present either in the same polymer or in additives and impurities. This absorption causes radical chain reaction mechanism initiated by carbon radicals due to the lack of oxygen. These primary radicals initiate and propagate the subsequent radical chain reaction and hence contribute to the degradation of materials[35]. However, Tmax of the UV irradiated BMU133/T3, BMU133/T2 are slightly increased compared to UV-BMU133. This can be ascribed to the protective effect offered by tracer particles due to their UV-screening properties inducing a good thermal stability

of the traced polymers towards ultraviolet radiations[27]. For visualization purposes, the TGA curves of UV-BMU133 and UV-BMU133-T3-0.1 after irradiation treatment are depicted in Figure 4.

3.3 Morphological properties It is commonly known that the degree of photo-oxidation decreases when moving deeper below the material surface owing to the limited penetration ability of UV light and oxygen diffusion. Therefore, the damage caused by UV irradiation is the formation of surface cracks inducing thereafter the embrittlement of the material[19,36]. Figure 5 compares the surface topography of the UV-BMU133 and UV-BMU133-T3-0.1 traced blend respectively. In both cases, cracks caused by UV irradiation can be clearly observed. However, the cracks on the surface of UV-PP are very coarse and propagate toward the core of the specimen while that for UV-BMU133-T3-0.1 are much finer. This suggests that damages caused by UV irradiation to the virgin BMU133 are more intense than those to the traced blend with T3 (UV-BMU133-T3-0.1). It can also be easily observed that the surface cracks are significantly reduced after addition of the tracer particles T3. The same type of surface cracks was observed for BMU133 filled with T2 tracer. As for T3 tracer, addition of the T2 or T3 microparticles can reduce the damage of UV light and oxygen diffusion on PP polymer and strengthens the resistance of PP to photodegradation.

Table 1. Thermal and crystallization data obtained from DSC and TGA analyses before and after UV irradiation.

Screw speed N=800 rpm

Samples BMU133-reference (B) UV-BMU133 BMU133-T3-0.1 UV-BMU133-T3-0.1 BMU133-T2- 0.1 UV-BMU133-T2-0.1

∆Hf

ΔHc

T max

(°C)

(J/gPP )

(°C)

(J/gPP)

(%)

(°C)

168.2 159.6 167.1 154.1 167.2 151.5

57.2 65.4 55.2 63.1 57.1 64.2

124.0 112.8 124.0 113.4 124.0 114.5

57.1 67.3 62.4 65.4 63.2 66.1

24.5 36.3 26.5 32.1 27.3 30.7

485.2 460.3 471.1 468.8 488.5 466.3

Figure 4. TGA curves of the UV-BMU 133 and UV-BMU133-T3-0.1 after irradiation. Polímeros, 25(6), 515-522, 2015

519

Massardier, V., & Louizi, M.

Figure 5. SEM micrographs of (a) BMU133, (b, c) UV-BMU133, (d) UV-BMU133-T3-0.1.

3.4 Effect of UV-irradiation on the mechanical properties of traced blends Before UV irradiation treatment, BMU133 and its traced blends processed at 800 rpm show a ductile behaviour (Table 2). However, after UV treatment, an intense decrease of both modulus and elongation at break is observed (Figure 6). This means that the photodegradation weakens the material and induces a fundamental change in the behavior of the as-exposed samples blends from ductile to brittle material. This behaviour is in agreement with the chemi-crystallization caused by the degradation process of PP, which contributes to the formation of surface cracks due to the contraction of the surface layer[16,18]. Furthermore, in addition to chemi‑crystallization, crosslinks and chain scissions can lead to the embrittlement of materials and it is also noted that the accumulation of defects (carbonyl, hydroperoxyde groups, etc.) on the chains in the amorphous phase induces the formation of tensile residual stresses[37,38]. In addition, the presence of surface cracks, as seen previously on SEM micrographs (Figure 5), can be considered as one of the reasons for the embrittlement of the material, causing 520

Figure 6. Elongation at break of samples before UV irradiation.

and after

Polímeros , 25(6), 515-522, 2015

Photodegradation of a polypropylene filled with lanthanide complexes Table 2. Mechanical properties (E: Young Modulus, σb : strain at break, εb: elongation at break) of samples before and after UV irradiation.

BMU133-reference (B) UV-BMU133 BMU133-T3- 0.1 UV-BMU133-T3- 0.1 BMU133-T2- 0.1 UV-BMU133-T2- 0.1

N= 800 rpm

Screw speed

Samples

σb

εb

(MPa) 420±10 204±04 410±08 198±08 430±08 187±02

(MPa) 13±01 8±01 13±01 9±01 12±01 8±01

(%) 187±10 87±07 251±10 121±10 224±07 93±10

a serious deterioration in the mechanical properties of the products even after short term exposures[19]. Consequently, the presence of these cracks facilitates the failure.

4. Conclusions Results from this research show that the addition of tracer particles into a commercial PP formulation can impart an improvement in photodegradation resistance of PP to UV-irradiation. The tracers were stable after ageing tests. From FTIR analysis, it is observed that the intensity of carbonyl absorption for traced blends in the wavenumber region [1700-1800 cm–1] is reduced in presence of tracer particles compared to the unfilled material. This is ascribed to the protective effect offered by tracer particles due to their UV-screening properties which can retard the oxidation of PP. In addition, TGA analyses reveal that the traced blends can maintain better thermal properties, after irradiation. A significant increase of the crystallinity rate and a decrease of the melting temperature were detected for unfilled UV irradiated PP. This might result from chemi-crystallization which can occurr when chain entanglements are broken as a result of chain scissions. The freed segments produced by molecular chain scissions induce the formation of new crystalline structure in the amorphous zone, especially at high exposure temperature and thus provoke the growth of pre-existing crystals of PP. These changes in thermal behavior are less pronounced for traced blends. After UV irradiation treatment, surface cracks appear on the surface of unfilled and filled BMU133 with T3. The intense surface cracks observed on the unfilled material is due to the photodegradation on the specimen surface. However, what is interesting here is the reduction of the surface cracks extent when adding T2 or T3 tracers. These cracks in the surface layer of degraded PP are observed by scanning electron microscopy (SEM). Furthermore, the mechanical tests confirm this finding and indicate a fundamental change in the behavior of the as-exposed blends. Such an effect is ascribed to chain scissions, accumulation of defects (carbonyls, etc.) and formation of surface cracks caused by the degradation process of PP. To sum up, the addition of lanthanide complexes to polymers is of importance as i) it allows their rapid identification by UV fluorescence spectrometry in order to facilitate the sorting and recycling of end-of-life products[1] and ii) limit the degradation of polymers under UV radiations. Thus, increased durability of polymer materials can be obtained by addition of lanthanide complexes such as T2 and T3. Polímeros, 25(6), 515-522, 2015

Impact Strength (kJ/m2) 7±01 3±01 12±01 5±01 11±02 6±01

5. Acknowledgements The authors would like to thank the National Research Agency (ANR) for its contribution to the funding of this work and for providing industrial orientations and scientific supervision to the research. Authors also wish to acknowledge P. Alcouffe and the “Centre de Microstructures et d’analyses, plateforme Lyon 1” of the University Lyon1 for his assistance.

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http://dx.doi.org/10.1590/0104-1428.1938

Thermal Stabilization study of polyacrylonitrile fiber obtained by extrusion Robson Fleming Ribeiro1*, Luiz Claudio Pardini2, Nilton Pereira Alves3 and Carlos Alberto Rios Brito Júnior4 Departamento de Química, Instituto Tecnológico de Aeronáutica – ITA, São José dos Campos, SP, Brazil 2 Departamento de Ciência e Tecnologia Aeroespacial, Instituto de Aeronáutica e Espaço, Instituto Tecnológico de Aeronáutica – ITA, São José dos Campos, SP, Brazil 3 Quimlab Científica Ltda, Jacareí, SP, Brazil 4 Universidade Federal do Maranhão – UFMA, São Luís, MA, Brazil

*robsonfleming@gmail.com

Abstract A low cost and environmental friendly extrusion process of the Polyacrylonitrile (PAN) polymer was viabilized by using the 1,2,3-propanetriol (glycerol) as a plasticizer. The characterization of the fibers obtained by this process was the object of study in the present work. The PAN fibers were heat treated in the range of 200 °C to 300 °C, which is the temperature range related to the stabilization/oxidation step. This is a limiting phase during the carbon fiber processing. The characterization of the fibers was made using infrared spectroscopy, thermal analysis and microscopy. TGA revealed that the degradation of the extruded PAN co-VA fibers between 250 °C and 350 °C, corresponded to a 9% weight loss to samples analyzed under oxidizing atmosphere and 18% when the samples were analyzed under inert atmosphere. DSC showed that the exothermic reactions on the extruded PAN co-VA fibers under oxidizing synthetic air was broader and the cyclization started at a lower temperature compared under inert atmosphere. Furthermore, FT-IR analysis correlated with thermal anlysis showed that the stabilization/oxidation process of the extruded PAN fiber were coherent with other works that used PAN fibers obtained by other spinning processes. Keywords: polyacrylonitrile fibers, extrusion process, stabilization process, carbon fibers.

1. Introduction Polyacrylonitrile (PAN) fibers have been the most used precursor for carbon fibers manufacture. Nowadays, PAN fiber represents about 90% of worldwide carbon fiber production. The production of PAN fibers, commercially in the form of copolymers, for over 50 years has been performed by wet spinning and dry spinning processes[1,2]. These processes are based on the use of organic solvents that dissolve the PAN polymer and a solution is formed. The solution is then pumped through a spinneret having a multiplicity of tiny holes[1,2]. The solution used in a wet spinning process usually consists of 10 to 30% by weight of PAN dissolved in a polar solvent, such as dimethylformamide (DMF), dimethysulfoxide (DMSO) or dimethylacetamide (DMAc)[3,4]. Although PAN fibers can be produced by either wet or dry spinning processes, wet spinning is used to produce nearly all commercial aerospace grade PAN-based carbon fibers. Due to the large amounts of solvents used and consequently environmental concerns, PAN industrial plants needs a solvent recovery plant. This contributes to an increase in PAN fiber final cost, which nowadays corresponds to nearly 50% of the carbon fiber final cost (~US$ 50.00/kg). So, carbon fiber uses are still concentrated only in areas where the importance of material performance overlaps the drawbacks of relative high cost production. However, the development of new processes without the use of toxic and high cost solvents, it could become an attractive technology route for PAN polymers. Several attempts have been made along the years to process PAN by using conventional thermoplastic

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extrusion technology. In this way, some processes were patented, mainly from BASF and Dupont, but they were not commercially viable because the need for water at high temperature (200 °C) as a melt carrier[5-7]. Besides, the extrusion process has to be accomplished at high pressure (3-7 MPa), and, as a consequence, robust equipment was needed for continuous extrusion. Others attempts have been made trying to obtain meltable PAN copolymers for fiber manufacture, which were later patented under trade names of Barex and Amilon[8,9]. In this case, the acrylonitrile polymerization was done by emulsion and the high cost of comonomers has threatened the economic viability of the products. Bortner[10], for instance, demonstrated successfully that using supercritical carbon dioxide as a plasticizer for PAN, which reduces the glass transition temperature and lead to a reduction in the PAN melting point, avoiding PAN degradation. In this case, PAN could be processed by injection, but the process has to be carried out at high pressure (6 MPa) with liquid CO2, in a pressure vessel (17.2 MPa), at 120 °C, to accomplish the saturation with acrylonitrile. PAN copolymers could also be processed by extrusion by using 1,2,3-propanetriol (glycerol) as a plasticizer[11]. The 1,2,3-propanetriol has a high boiling point (~290 °C) and exhibits high miscibility with PAN, due to the three hydroxyl groups in its backbone which provides a polar nature substance. Thus, these groups interact with nitrile group and can delay the PAN nitrile group’s cyclization[11].

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Ribeiro, R. F., Pardini, L. C., Alves, N. P., & Brito, C. A. R., Jr. Thus, the PAN polymer melt before crosslinking and suppresses the exothermic reactions related to the cyclization reactions which takes place over 200 °C. In that way, PAN fibers can be spun by using conventional extrusion process (melt spinning)[4,12]. In relation to acrylic fiber production this approach can lead to a substantial cost reduction in relation to the traditional wet and dry spinning processes[1]. Besides, it is environmentally friendly since avoids human exposure to toxic solvents. Also, the use of glycerol as PAN plasticizer on industrial scale becomes an alternative route for using glycerin produced by biodiesel industry waste. Conversion of PAN fibers to carbon fibers requires a critical thermal stabilization stage which has great influence on the final properties of the carbon fibers. The stabilization process is commonly performed in air between 200 and 300 °C[13,14]. During this process, PAN undergoes a number of physical and chemical changes which converts the linear PAN molecular chains to an aromatic ladder structure suitable for further heat treatment (carbonization) and conversion to carbon fibers. Although the chemistry of stabilization process is complex, this stage consists basically of exothermic chemical reactions, including cyclization of nitrile groups, dehydrogenation, oxidation and crosslinking among the polymer chains[15,16]. Several mechanisms have been proposed to explain the structure of PAN fiber backbone and its stabilization oxidative process. Fourier transform infrared (FT-IR) spectroscopy, microscopy and thermal analysis have proven to be excellent techniques to study the structural evolution of PAN fibers during the stabilization. In this work, PAN fibers spun by the extrusion process were submitted to heat treatment processes, related to stabilization stage. The main functional groups and thermal behavior of extruder PAN fibers were identified and characterized. Results were compared with the PAN fibers spun from conventional solvent wet and dry spinning processes.

2. Experimental 2.1 Materials Extruded PAN fibers were obtained by continuous extrusion process. In order to obtain PAN fibers, a PAN copolymer having 94%/mass of acrylonitrile monomer (AN) and 6%/mass of vinyl acetate (VA) was used. The PAN co-VA was provided by Radicifibras (Brazil/SP) having a number average molecular weight (Mn) of 45000 g/mol. The glycerol was provided by Vetec Química (Brazil) having 99.5% purity, Mn of 92 g/mol, melting point at ~18 °C and boiling point at ~290 °C.

2.2 Melt spinning process for PAN fibers The process of melt spinning of PAN co-VA fibers is based in two stages. The first step corresponds to the formation of pellets of PAN plasticized with glycerol by a single screw extrusion and formation of PAN fiber by cascade extruder equipment. The compound is prepared by mixing 70%/mass of PAN co-VA powder, 28%/mass of glycerol and 2%/mass of glycol additives[11]. The pre‑extrusion compound is homogenized and put into the hopper extruder equipment. Figure 1 shows a schematic view for PAN fibers extrusion. Afterwards, the compound melts at 200 and 220 °C and it is transported by a single screw through the barrel. The molten polymer is compressed up through a circular die duct of 6 mm diameter and the compound has been conformed into filaments. The filaments are then pulled and stretched reducing its cross section to a diameter of up to 3 mm. Further, the 3 mm filaments are cut by a rotary knife turning it into pellets of plasticized PAN. The second step is concerned with extrusion of PAN pellets for obtaining fibers. The Figure 2 shows a schematic view of the cascade extruder equipment used in this stage. Firstly, PAN co-VA copolymer is fed, in the form of pellets directly into the hopper. The pellets undergo melting and begin flow in the screw extruder leaving the exit gate.

Figure 1. Schematic view of a single screw extruder for obtaining pellets of plasticized PAN. 524

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Thermal Stabilization study of polyacrylonitrile fiber obtained by extrusion Afterwards, the molten PAN die drops by gravity into the cavity of a secondary extruder. The release of undesirable volatiles occurs during this event, avoiding wasting material, as occurred in the degassing of the single screw type extruder. Due to the strong backpressure generated in the die output, part of PAN melting run off by degassing orifice (area of lower pressure). The die output of spinning is coupled to the secondary extruder and the filaments emerging from the spinneret are directed by a guide cylinder to the roller guides. A hot air box chamber, working at 170 °C, is positioned just after the spinning die and solidification of PAN fiber takes place by fast cooling.

2.3 Thermal treatment Thermal stabilization of the extruded PAN fibers was carried out in a furnace under stagnated air atmosphere. No stretching was applied to the fiber material. The treatment temperature were set at 200 °C, 220 °C, 240 °C and 260 °C. At each temperature, an isotherm treatment was performed at 30 min, 60 min and 90 min. Heat treatments above 260 °C were not possible to be done, because the fibers showed abrupt shrinkage leading to rupture.

2.4 Characterization FT-IR spectra of untreated PAN fibers were obtained on the Perkin Elmer, Spotlight 400 FTIR Imaging System model spectrometer. For the heat treated PAN fibers a FT‑IR Pike Technologies Varian 640-IR model spectrometer was used. Both analyses were performed by ATR method in 4000-560 cm–1 range at 4 cm–1 resolution and 10 scans. Cross section fracture morphologies from the extruded PAN fibers were obtained by using LEO 435 Vpi scanning electron microscope. The samples were vacuum metalized with a thin film of gold-palladium alloy. The analysis was performed using 20 kV. Thermogravimetric Analysis (TGA) and Differential Scanning Calorimetry (DSC) were performed to evaluate the thermal behavior of the extruded PAN fibers. TGA and DSC analysis were performed in a Shimadzu equipment model. The samples were placed in aluminum pan, with mass of 6.5 mg. The samples were heated from ambient temperature up to 350 °C, at a heating rate of 5 °C/min. All the scans were performed in duplicate, under synthetic air and nitrogen flowing at 100 mL/min.

3. Results and Discussion 3.1 FT-IR analysis

Figure 2. Schematic view of the cascade extruder for PAN fibers.

The FT-IR spectra of the extruded PAN fibers are show in Figure 3a. The main chemical group assignments for the extruded PAN fibers were identify and have the same polymeric structure of the PAN fibers obtained by wet and dry spinning process. Of particular interest is the assignment at 2240 cm-1 which is attributed to the C≡N stretching of acrylonitrile unit in the polymer chain and the bands at 2940 cm-1 (νC−H in CH2), 1452 cm-1 (δC−H in CH2) and 1370 cm-1 (δC−H in CH), which are characteristics of aliphatic CH groups along the PAN backbone[17,18]. The weak absorption at 1626 cm-1 may be attributed to C=C (νC=C)[19]. A strong assignment appears at 1738 cm-1 which is attributed to the carbonyl group (C=O) stretching, which is due to the

Figure 3. FT-IR spectra of (a) extruded PAN co-VA fibers and (b) commercial PAN co VA fibers obtained by wet spinning process having composition of 94%/mass of AN and 6%/mass of VA. Polímeros, 25(6), 523-530, 2015

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Ribeiro, R. F., Pardini, L. C., Alves, N. P., & Brito, C. A. R., Jr. ester and acids comonomers used in the polymerization of the PAN polymer. Ouyang[17], Mathur[20] and Wangxi[21], also found an evident assignment of the carbonyl group of great intensity in PAN fiber copolymerized with itaconic acid and vinyl acetate, using IR technique. Besides the carbonyl band, an intense assignment at 1238 cm-1 (νC−O) is found which is related to the acetate group due to the vinyl acetate copolymer used in the synthesis of the PAN co-AV. At 1042 cm-1 another strong band appears assigned to the 1,2,3-propanetriol, used as a plasticizer during the PAN fiber spinning. This assignment band characterizes the C–O groups, linked to the primary carbons of the 1,2,3-propanetriol. The broad band around 3500 cm-1 is a characteristic band of O–H vibrations, which can be assigned to O–H of the 1,2,3-propanetriol, as well as for O–H from absorbed water in the fiber surface. For comparison, a similar FT-IR ATR spectrum was obtained from a commercial PAN fiber having the same copolymer composition (94% AN/6% VA), but using the process of wet spinning with DMF solvent. It can be seen from Figure 3b, that the bands at 2936 cm-1 (νC−H in CH2), 1452 cm-1 (δC−H in CH2), 1362 cm-1 (δC−H in CH), 2244 cm-1 (νC≡N), 1626 cm-1 (νC=C) and 1732 cm-1 (νC=O), which are characteristic bands of PAN co-VA backbone, are similar to

the assignment bands obtained for PAN fibers by obtained by the extrusion process showed in Figure 3a.

3.2 Fracture morphology of PAN co-VA fibers The fundamental fiber structure needed to develop high strength must be created during the initial fiber formation step. The Figures 4 and 5 exhibits the SEM images of cross section and longitudinal section of the extruded PAN co‑VA fibers, respectively. It can be observed from Figure 4 that the extruded PAN co-VA fibers exhibit a regular round cross section. This circular cross section is a result from the melted PAN co-VA copolymer extruded through the spinneret capillaries. A regular round cross section is the most appropriate morphology for the thermal conversion process to high performance carbon fibers, because it provides a homogeneous shrinkage and mass reduction, keeping other process parameter under control, throughout the fiber during the thermal treatment cycle. Regular round cross section and strict heat treatment process control reduces the probability of defects that may affect the final mechanical properties of the carbon fibers. The magnified image from Figure 4 shows the cross section of the PAN co-VA where voids can be seen, as

Figure 4. Cross section SEM images of fracture morphologies of extruded PAN co-VA fibers.

Figure 5. Longitudinal section SEM morphology images of extruded PAN co-VA fibers, showing longitudinal polymer fibrils (a) and polymer granules at the core of the fiber (b). 526

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Thermal Stabilization study of polyacrylonitrile fiber obtained by extrusion indicated by the red circle. In this case, the release of plasticizer during extrusion can be the source of such defects. Voids are harmful and undesirable since they compromise the tensile fracture of the precursor fiber. According to Ji[23], voids are stress concentrators and a source of the tensile fracture. Efforts in the process to obtain carbon fibers are addressed to reduce the number and the size of voids, and so a high tensile strength can be obtained. Figure 5a shows a longitudinal surface view of the extruded PAN co-VA fibers. The polymer fibrils are oriented preferably in the direction parallel to the longitudinal axis of the fiber. This is due to the fact that extrusion process orients the molecular chains of the PAN co-VA polymer longitudinally during spinning even before the drawing step. The image from Figure 5a shows that granules showed in Figure 5b are the fracture ends of fibrils and the transverse section of these granules are randomly distributed. Chain‑like molecules present when PAN polymer are spun is a well‑known morphology present in fiber precursors for carbon fibers, which indicates an adequate orientation in the axial direction. Besides, a random distribution of the fiber ends granules in the transverse direction results in a random transverse texture[22,23].

3.3 Thermogravimetric Analysis (TGA) The evolution of weight loss of the extruded PAN co‑VA fibers was measured as a function of temperature under oxidative atmosphere and under inert atmosphere, as shown in Figures 6a and 6b, respectively. Main oxidative reactions from the thermal stabilization stage of these fibers takes place up to 350 °C[2,24]. The curves shown in Figure 6 shows that extruded PAN co-VA fibers exhibits three characteristic regions of weight loss. The weight loss of PAN co-VA fibers under oxidizing atmosphere shows a weight loss of ∼8%, and under inert atmosphere the weight loss is ∼ 6%, up to ∼110 °C. These levels of weight loss are attributed to adsorbed humidity, in addition to the release of 1,2,3-propanetriol, used as a plasticizer during fiber spinning. The presence of humidity and 1,2,3-propanetriol can also be demonstrated by free OH characteristic assignment detected in the FT-IR analysis, as showed the Figure 3a.

The second stage of weight loss is in the temperature range of 125 °C to 250 °C, which corresponds to a 15% weight loss for both samples analyzed under oxidizing atmosphere and inert atmosphere. In this case, weight loss can be attributed to dehydrogenation reactions and release of gases, such as CO2[25]. The third step, between 250 °C and 350 °C, corresponds to a 9% weight loss to samples analyzed under oxidizing atmosphere. On the other hand, when the samples were analyzed under inert atmosphere the weight loss (∼18%) is more significant compared to the analysis showed in Figure 6(a), which is approximately twice the weight loss measured under oxidizing atmosphere. Thus, the yield on carbon from extruded Pan co-VA fiber under inert atmosphere is smaller than the resulted yield measured under oxidizing atmosphere. Anyway, in both samples, the weight loss can be attributed to the release of volatiles, such as H2O, CO, CO2, CH4, NH3 and HCN[25,26]. As a matter of fact, under oxidizing atmosphere, the extruded PAN co-VA fibers exhibit a gradual weight loss. From 250 °C to 350 °C the oxidation reactions occur with greater intensity and acting as a dehydrogenation agent during the conversion of C—C bonds to C=C bonds. Oxygen molecules also generate oxygen-bearing groups in the polymer backbone, such as OH, C=O and COOH. These groups promote intermolecular crosslinking of the polymer chains and provide greater stability to sustain high temperature carbonization treatment[17,18]. The formations of ammonia and hydrogen cyanide have been reported by several workers but nobody has suggested a detailed pathway for these reactions, especially in the absence of oxygen[27]. According to Xue[27], the formation of NH3 is assumed to be due to the terminal imine group from the cyclized structure and the formation of HCN is assumed to be due to the uncyclized nitrile group. The analysis carried out in the extruded PAN co-VA fiber under inert atmosphere favors a loss in properties because a non-stable structure is prone to oxygen attack. The yield of carbon after heat treatment schedule of a precursor fiber is a key issue in the carbon fiber manufacture. If the carbon yield is low, the tensile strength and elastic modulus will tend to exhibit lower values[2]. This means that heat treatment

Figure 6. TGA curve of the extruder PAN co-VA fibers heated at 5 °C/min under (a) oxidizing atmosphere and (b) inert atmosphere. Polímeros, 25(6), 523-530, 2015

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Ribeiro, R. F., Pardini, L. C., Alves, N. P., & Brito, C. A. R., Jr. of PAN co-VA under oxidizing atmosphere is convenient during the stabilization process in order to get carbon fibers with better properties.

3.4 Differential Scanning Calorimetry (DSC) The Figure 7 shows a DSC analysis where an exothermic thermal event occurs, related to the extruded PAN co-VA fibers stabilization stage, under oxidizing synthetic air atmosphere and inert atmosphere. In general, exothermic events occurring on polyacrylonitrile polymer, takes place owing to the cyclization reactions of nitrile groups, dehydrogenation, crosslinking and oxidative reactions, when treatment is performed at oxidizing atmosphere[14, 19]. The exothermic reactions occurring on the extruded PAN co-VA fibers initiate at ∼190 °C (Ti), goes to a peak at 290 °C and ended at ∼340 °C (Tf). The extent of the stabilization event (ΔT = Tf – Ti), under oxidizing atmosphere, is around 150 °C and releases an amount of energy of about 93 J/g. Under inert nitrogen flow atmosphere the exothermic reactions initiate about 195 °C (Ti), having a peak at 297 °C and ended about 318 °C (Tf). The extent of the stabilization

Figure 7. DSC curves of extruder PAN co-VA fibers heated 5 °C/min under synthetic air flow and nitrogen flow.

event (ΔT = Tf – Ti), under inert atmosphere, is around 123 °C and releases an amount of energy about 70 J/g. It is observed that under inert nitrogen atmosphere the reaction exhibits a short extension (ΔT = Tf – Ti) and peaks of higher magnitude and narrower are present. Thus, under inert atmosphere, the exothermic reactions release energy in a shorter time interval, the weight loss rate is increased and further reaction may become uncontrolled. This can be attributed to the lack of crosslinking generated by oxygen groups in the polymer backbone. DSC results showed that extruded PAN co-VA fibers analysed under oxidizing and inert atmosphere have a main thermal event in the temperature range of 180 °C to 350 °C, which is in accordance with the temperature range for exothermic reactions related to the thermal stabilization step of PAN based carbon fibers. Other studies that evaluated PAN fibers spun by the wet and dry spinning process, using solvents such as DMSO and DMF, also found similar results [28, 29].

3.5 FT-IR analysis of the PAN co-VA fibers heat treated at different temperatures and times Figures 8 and 9 shows FT-IR spectra from extruded PAN co-VA fibers heat treated at different temperatures and times. It is observed, as shown in Figure 8a, that the PAN fibers heat treated at 200 °C for 30 min shows peak assignments similar to the spectrum of the untreated PAN fibers. However, an increase in the dwell time at 200 °C, leads the weak band at 1631 cm-1, assigned to the stretching C═C, being moved for a new assignment band around 1595 cm-1. According to Rahaman[13] and Ouyang[17], this weak band at 1595 cm-1 seems to be due to a combination of vibrations of C═C and C═N stretching, and ═N—H in-plane bending of the ladder frame structure of the stabilized PAN. This results in the progress of cyclization and dehydrogenation of PAN fibers, as can be confirmed by DSC curve (Figure 7), which shows at 200 °C the beginning of the thermal event. Heat treatment of PAN co-VA fiber at 220 °C shows a characteristic assignment band, related to the cyclization, between 1590 cm–1 and 1596 cm–1 which is more intense.

Figure 8. FT-IR spectra of extruder PAN co-VA fibers heated at (a) 200 °C and (b) 220 °C, for different times. 528

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Thermal Stabilization study of polyacrylonitrile fiber obtained by extrusion

Figure 9. FT-IR spectra of extruder PAN co-VA fibers heated at (a) 240 °C and (b) 260 °C, for different times.

After 90 min treatment at 220 °C, as shown in Figure 8b, the cyclization assignment band (1590 cm-1) finally overtakes the 2242 cm–1 band. It is observed, according to Figure 9a, that after 30 min treatment at 240 °C, the 2244 cm–1 assignment band related to the nitrile groups shows a significant decrease in intensity. Concomitantly, the band corresponding to cyclization (1590 cm–1) continues to increase, becoming very strong. Besides, it is interesting to observe that the assignment band at 1450 cm–1, which is attributed to the C—H in (–CH2) bending, begins to decrease in intensity becoming a shoulder of the 1365 cm–1 band, which assigned to the C—H in (–CH) bending. This result shows that during the dehydrogenation reactions, the hydrogen atoms linked in (–CH2) are broken more frequently. When heat treatment of extruded PAN co-VA fibers are performed at 260 °C, a weak band appears around 2200 cm–1. This band was suggested to be assigned to the unsaturated nitrile groups arose from dehydrogenation, or from tautomerization of the ladder polymer[17,19]. The PAN fibers heat treatment at 260 °C for 90 min still shows a weak band at 2240 cm-1 which indicates that a small fraction of nitrile groups fails to complete the cyclization. According to Mathur[20], the characteristic band that gives evidence of the presence of the vinyl acetate unit in the AN/VA copolymer is the carbonyl band, around 1730 cm–1. This band appears with great intensity in the untreated PAN co-VA fibers, studied in the present work. But, as shown in Figure 9a, this band begins to decrease in intensity after 30 min at 240 °C, and almost disappears after 60 min of treatment. In the spectrum of the Figure 9b the carbonyl band is vanished. These results show that the oxygen did not generate groups as C═O in the polymer back-bone.

4. Conclusions PAN co-VA fiber was obtained by conventional extruder process. This process (melt spinning) converted a pure precursor directly into fiber form and not involved the added expense of solvent. Polímeros, 25(6), 523-530, 2015

The FT-IR analysis results showed that PAN co-VA fibers, obtained by extrusion process, exhibited the main functional groups, similar to the ones founded from other PAN fibers obtained by wet and dry spinning. The thermal analysis (TGA and DSC) showed that stabilization stage of the PAN co-VA fibers beyond to being consistent with other fibers obtained by wet and dry spinning process, when it was performed in an oxidizing atmosphere showed a more stable release of volatiles, allowing a higher yield of carbon.

5. Acknowledgements We thank the Quimlab Company for the viability of the present work and CNPq for the financial support.

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http://dx.doi.org/10.1590/0104-1428.1732

The influence of long chain branches of LLDPE on processability and physical properties Paula Cristina Dartora1, Ruth Marlene Campomanes Santana1* and Ana Cristina Fontes Moreira2 Department of Materials Engineering, Universidade Federal do Rio Grande do Sul – UFRGS, Porto Alegre, RS, Brazil 2 Politecnic Institut, Universidade Estadual do Rio de Janeiro – UERJ, Nova Friburgo, RJ, Brazil 1

*ruth.santana@ufrgs.br

Abstract Two polyethylene-based on single-site metallocene catalyst (mLLDPE) were selected to characterize the effect of long chain branching (LCB) on blown film processability, optical and mechanical properties. 13C NMR and parallel plate rheology were used to identify LCB presence on LLDPEs. Blown films were produced from 100% LLDPEs using three different machine direction (MD) stretch ratios. When the same processing conditions for the two LLDPEs grades were used, better processability was observed for LLDPE with LCB. In relation to mechanical and physical properties, Elmendorf tear and optical properties were highly influenced by the presence of LCB. Tear resistance is affected by film orientation and is inversely proportional to the level of LCB in the polymer. It was observed a reduction of 50% in the MD tear strength when comparing with the polymer without LCB. However, haze decreases significantly with the presence of LCB, about 40%. Keywords: LLDPE, metallocene, long chain branching, NMR, rheology, mechanical properties.

1. Introduction Polyethylene is the most useful polyolefin in the world. It is available commercially as groups of polyethylene: high density polyethylene (HDPE), low density polyethylene (LDPE), linear low density polyethylene (LLDPE), ultra high molecular weight polyethylene (UHMWPE) and very low density polyethylene (VLDPE)[1]. Flexible packaging marketing is the major application for LLDPE. Packaging fulfills four functions: containment, protection, convenience and communication[2]. For communication one can understand that the package usually sports the name of the product and nutritional information, for example. But the package also shows the product to the costumer, so it is important that the film have low haze[2,3]. They are basically produced by blown and cast film processes. Structural parameters, such as density/crystallinity, molecular weight and its distribution, short chain branching (SCB) / long chain branching (LCB) length and amount and crystalline morphology are the key factors that control the properties. LLDPE are produced by copolymerization between ethylene and an alfa-olefin comonomer such as 1-butene, 1-hexene or 1-octene. It results in an ethylene/alfa-olefin copolymer with many short chain branches along the polymer backbone. Ultimate developments in metallocene catalysts allowed adding LCB on LLDPE structure during copolymerization[4]. The mechanism of LCB formation is not well known, but the most accepted is a random intermolecular reaction although in some cases this mechanism does not explain the phenomena observed[4,5]. Another way to obtain LLDPE with long chain branches is mixing LLDPE with peroxides, using peroxide concentrations below the critical gel formation concentration[6].

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There remains a need to find sensitive methods to characterize LCB architecture[7]. Gel permeation chromatography (GPC) was the first method used to detect LCB. However, it was necessary to fractionate the polymer before characterization and it was not possible to obtain quantitative information[8,9]. To quantify LCB levels, more complex analyses are needed. One widely used technique is nuclear magnetic resonance (NMR)[10]. This technique, however, has a limitation. When the alfa-olefin comonomer used to produce the polymer has more than six carbon atoms, the chemical shift observed for LCB is the same observed for the comonomer[10]. Recently, studies showed that it is possible to distinguish branches longer than six carbons, up to twenty carbon atoms, but it is necessary high resolution[11]. Nevertheless, the most sensitive method is the rheological method. Small levels of LCB, such as 1LCB / 10.000C affect the rheological behavior of the polymer[6,12]. LCB has distinct effects on different rheological quantities (zero shear-rate viscosity - η0 - and strain hardening). This is caused by the fundamental differences between the molecular mobility of linear and long branched chains. Strain hardening and η0 increase when LLDPE presents LCB[13]. To be able to use rheological data to compare two or more polyethylene grades, it is necessary that they have the same polydispersity[14], because long molecules in a narrow molar mass distribution can create the same rheological behavior in elongation as long-chain branches[13]. One way to detect LCB using rheological data is considering the flow activation energy of the LLDPE. LLDPE usually has a flow activation energy smaller than 28 kJ/mol. On the other hand, LDPE, which is a highly branched polymer and presents a flow activation energy

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Dartora, P. C., Santana, R. M. C., & Moreira, A. C. F. around 60 kJ/mol. When LLDPE resins have LCB, the flow activation energy increases, reaching values of about 45 kJ/mol[14]. LCB is commonly introduced into the fractions with higher molecular weight[15], decreasing the viscosity of these fractions[8]. The decreasing of viscosity improves the processability of the polymer[7,8,14]. It is especially important when metallocene-based polymers are considered, as these polymers have problems with processability caused by their narrow molecular weight distribution[14]. SCB and LCB affect polymer’s density, crystallinity, resistance and processability, as they change the polymer’s structure[15]. It is recognized in literature that higher levels of LCB in blown film resins improve bubble stability[7]. It is also known that LCB governs die swell, melt strength, and environmental stress crack resistance in blow molding operations, orientation in film, sag resistance in pipe and geomembranes and shear thinning and melt fracture in all extrusion processes[4]. Mechanical and physical properties are affected by polymer structure[16]. For example, haze is affected by polymer crystallinity and/or crystal structure, surface imperfections and bubbles or particles (additives) in the film[17]. Recent studies show that the LCB content in polymer chain affects the polymer crystallinity, crystal thickness, tensile strength, tensile modulus and rheological characteristics of mPELBD (m means polymer prepared via metallocene catalyst)[18]. Since LCB affects polymer crystallinity and crystal structure, it consequently affects film haze[19]. Elmendorf tear resistance is strongly dependent on film orientation. When a film is produced with an LDPE resin, increasing the film orientation increases Elmendorf tear resistance on machine direction (MD). However, when the raw material is a LLDPE, the opposite behavior is observed, decreasing significantly MD Elmendorf tear resistance when the film orientation is increased[20]. Structures of polyethylene blown films have been studied for a long time, but some concerns and controversy still exist, and some structural features and physical behaviors are not completed understood. The aim of the present work is at evaluating the effects of LCB on the blown film properties of 1-hexene-based LLDPE resins, in order to further define and understand the processing-structure-property behavior of these resins.

2. Experimental 2.1 Materials Two mLLDPE grades from Braskem were used in this work. Metallocene grades were produced with different types of catalysts but all of them used 1-hexene as comonomer. Table 1 shows details about the resins used. It is important to say that the designation of the catalyst as metallocene type A or B was only used to show that the resins were produced with different catalyst systems. 532

Table 1. Basic resin characteristics. Samples Density (g/cc) Melt flow index (g/10min)a Catalyst Comonomer (%wt)b Mw × 10–3 (kg/mol)c Mw/Mnc

mLLDPE A 0.9176 1.02

mLLDPE B 0.9194 0.58

Metallocene type A 8 123 1.96

Metallocene type B 8 115 3.39

190°C/2.16kg. bComonomer content obtained from 13C NMR. cMw and MWD were obtained from GPC. a

2.2 Resins characterization Samples were analyzed by 13C NMR on a Varian Wibe bore 400, with a 5mm probe. Samples were prepared by dissolving 50 mg of polymer in 0.7 mL of ortodichlorobenzene and 0.2 mL of tetrachloroethane-d2. For long chain branching (LCB) quantification, Equation 1 was used, where α is the medium’s intensity of LCB carbon atoms and Ttot is total carbon intensity[10]. (1/ 3) α /  4 Branches /10000 = carbons   ×10 (1) (TTot ) 

The samples were also analyzed by MCR501 Physica Anton Paar rheometer. Samples were prepared by compression molding, in circular shape, with 2.5 cm of diameter and 2 mm of thickness. A tension of 200 Pa was applied on a frequency sweep mode, from 0.001 to 100 Hz, at 190, 200 and 210 °C. Equation 2 was used to calculate flow activation energy for all samples[21]. η= B × e(

− Ea / ( R.T ) )

(2)

Where: η = apparent viscosity (Pa s); B = pre-exponential factor; Ea = flow activation energy (J mol–1); R = the universal gas constant (8.314 J mol–1 K–1); T is absolute temperature (K).

2.3 Films production All the blown film samples (100% LLDPE) were made on a Carnevalli CHD60 blown film line using typical linear low-density (LLDPE) conditions as follows: 200 mm die diameter; 1.8 mm die gap; 800 rpm screw speed; 2.2:1 blow up ratio (BUR); freeze line high (FLH) of about 60 cm, temperature profile from 180 to 200oC and three different film thickness: 35 µm, 60 µm and 100 µm. For output analysis, two indexes were observed using the 35 µm‑thickness samples.

(

)

Energy index = amperage / output A h kg −1

(

(3)

)

Output index = melt pressure / output bar h kg −1 (4)

These indexes represent two important aspects for plastic industry: energy consumption and limit of production. Energy consumption is a cost indicator; it means that with lower Polímeros , 25(6), 531-539, 2015

The influence of long chain branches of LLDPE on processability and physical properties energy consumption, it is possible to reduce cost per kg of production. Output index is related to capacity of motor load; in other words, extruders have a limit of melting pressure. This parameter depends of the type of material; some of them flow easily and as a result the melting pressure is low. Other polymers are difficult to flow, increasing the melting pressure leading to stop the production. The film samples will be identified in this paper according to Table 2. Samples mLLDPE A_35 and mLLDPE B_35 were chosen to perform bubble stability studies, in which the only variable was the screw speed. The screw speed was increased gradually up to 75% of the machine capacity (1450 rpm) to evaluate the bubble stability of both materials. Films with 35 μm were chosen because they are the most produced by packaging industry and they have the most critical processing conditions (higher MD stretch ratio).

2.4 Film characterization The Elmendorf tear properties of all blown films were measured according to ASTM D-1922, using a TMI Universal Tear Tester. Haze analysis follows the ASTM D1003 standard method. All samples were analyzed by a BYK-Gardner equipment. Gloss analysis at an angle of 45° was carried out following ASTM D2457 also using BYK-Gardner Table 2. Samples identification. Sample mLLDPE A_35 mLLDPE B_35 mLLDPE A_60 mLLDPE B_60 mLLDPE A_100 mLLDPE B_100

Film Thickness (µm) 35 35 60 60 100 100

MD stretch ratio 115 115 67 67 40 40

equipment. For the three analyses aforementioned, 10 test specimens were used. Film crystallinity was determined by DSC analysis, with a heating rate of 10 °C/min, from –20 to 200 °C. The reference value for PE 100% crystalline used was 286.6 J/g[22].

3. Results and Discussion 3.1 Resins characterization It is possible to observe in Table 1 that the resins have the same density and amount of comonomer, but they have a different molecular weight distribution (Mw/Mn). It can be assumed that they are similar in their molecular structures in relation to SCB amount and molecular weights[23]. Initially the samples were characterized using 13C NMR and rheological test. Comparative NMR spectra can be observed on Figure 1. The chemical shifts usually observed when ethylene is polymerized using 1-hexene as comonomer are present on both spectra, according to Table 3[10]. Two characteristic LCB chemical shifts (32.32 ppm and 22.92 ppm), however, are present only for mLLDPE B. These two peaks were observed only for sample mLLDPE B and correspond to the insertion of a branch with six or more carbon atoms. As the polymer was produced using 1-hexene as comonomer, the branching formed by the comonomer insertion could not have more than four carbon atoms. The LCB level on sample mLLDPE B is about 4.7 LCB/10000C. To obtain the level of LCB, Equation 1 was used[10]. Rheological analyses of the samples were carried out at three different temperatures, as shown in Figure 2. It can be observed that the rheological behavior of both samples is different. Comparing both figures, it can be observed that the sample mLLDPE A (Figure 2a) presents the Newtonian plateau for complex viscosity when low frequencies are

Figure 1. 13C NMR spectra for LLDPE resins. Polímeros, 25(6), 531-539, 2015

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Dartora, P. C., Santana, R. M. C., & Moreira, A. C. F.

Figure 2. Rheological data obtained for the sample (a) mLLDPE A; (b) mLLDPE B.

applied (< 1s–1), while the sample mLLDPE B does not present the Newtonian plateau. These differences can be explained in accordance to earlier studies[24-27], where it is shown that the broad Newtonian plateau on the complex viscosity is characteristic of polymers that present a narrow molecular mass distribution. These are in agreement with the data shown on Table 1, where it can be seen that the mLLDPE A sample has a narrower molecular weight distribution than mLLDPE B[3,4]. It is observed that the complex viscosity of mLLDPE A is lower than the complex viscosity of mLLDPE B considering the same temperature and low frequency, although both present Mw of the same order of magnitude (Table 1). This behavior can be attributed to the larger entanglement of the mLLDPE B, caused by the presence of LCB; so, the entanglement velocity is higher than the disentanglement velocity. On the other hand, when high frequencies are applied, the situation is reversed and mLLDPE B’s viscosity decreases quickly, becoming smaller than mLLDPE A’s viscosity after ~500 s–1[24]. LCB affects the viscosity of polymers in two ways: 1) the polymer with LCB has higher molecular weight entanglement compared to its linear polymer of the same Mw and same chemical structure; 2) the disentanglement of branched polymer is easier compared to linear polymer under shear force/stress[18]. The higher complex viscosity values observed on mLLDPE B are typical of LDPE, which are highly branched, and can indicate the presence of LCB[13]. When long branches are present in the polymer, the entanglements increase and consequently the complex viscosity increases as well as the molecular mobility is reduced[6,24]. When the rheological characterization was carried out with higher temperatures, the complex viscosity showed smaller values for both samples, as expected[24]. Table 4 shows data obtained for activation energy from Equation 2. The mLLDPE A flow activation energy is at the same level of classical short chain branched (SCB) LLDPE resins. On the other hand, the mLLDPE B flow activation energy is close to the corresponding energies of LDPE resins[14]. As 13C NMR data, flow activation energy calculated from rheology data showed an indicative of LCB only for mLLDPE B[14]. Thereby, NMR and rheology analyses confirmed that only one of the samples presents LCB. 534

Table 3. Chemical shifts typically observed for ethylene/1-hexene copolymer with the corresponding carbon and sequence assignments. Chemical Shift (ppm) TCE(d2)a

Carbon Assignment

Sequence Assignment

38.14

Methine

EHE

34.55

αδ+

EHEE+EEHE

34.17

4B4

EHE

30.94

γγ

HEEH

30.49

γδ+

HEEE+EEEH

30.00

δ+δ+

(EEE)n

29.56

3B4

EHE

29.38

3B4

EHH

27.28

βδ+

EHEE+EEHE

27.08

βδ

HHEE+EEHH

23.42

2B4

EHE

14.28

Methyl

EHE

Deuterated Tetrachloroethane (solvent); H = Hexene; E = Ethene.

Table 4. Flow activation energy. Samples mLLDPE A mLLDPE B

Flow activation energy (kJ/mol) 28 43

3.2 Films production Bubble stability and maximum output were evaluated during blown film extrusion for samples with 35 µm-thickness films. Both samples presented dimensional stability even at high output rates, of about 1450 rpm. However, it was not possible to maintain the usual bubble shape for sample mLLDPE A when the screw speed was higher than 1200 rpm. LLDPE films are usually produced with the bubble shape observed on Figure 3a. When the screw speed became higher than 1200 rpm, it was necessary to change the bubble shape to the conformation used for LDPE films production, shown on Figure 3b. The main goal of the blown film process is to manufacture a stable film with good physical and optical properties at a maximum production rate. In this case, the bubble shape is controlled in the area between the die exit and freeze Polímeros , 25(6), 531-539, 2015

The influence of long chain branches of LLDPE on processability and physical properties

Figure 3. (a) represents “LLDPE’s bubble shape”; (b) represents “LDPE’s bubble shape”.

line height, which leads to reduced product properties, line failures, and large amounts of film scrap. Such instabilities decrease significantly the window of stable processing conditions for blown film production[28]. The maximum flow rates measured for all samples were at the same range, close to 130 kg h–1. A high bubble stability, high production rate and better FLH control were obtained by mLLDPE B sample. In relation to processability, described here as energy and output indexes, data are shown on Table 5. These indexes are inversely proportional to energy saving and output rate. Also it is observed that mLLDPE B has lower energy and output indexes, which means that mLLDPE B needs lower levels of energy (–13%) and lower melt pressure (–16%) to produce the same weight of film (1 kg) in comparison with mLLDPE A. In other words, a customer using mLLDPE B resin is able to increase their productivity with a lower energy cost, so they became more competitive. Beyond processability characteristics, LLDPE resins have to present good optical and mechanical properties[28]. The positive effect of LCB on process parameters is already demonstrated and now, in addition, the LCB effect on optical and mechanical properties of the films will be evaluated.

3.3 Film characterization Table 6 shows haze and gloss properties for all samples. It is possible to note that mLLDPE A exhibits higher variations on the results. The probable cause of this is the stripes formed during blown film extrusion. Figure 4 shows these stripes for mLLDPE A and no stripes for mLLDPE B during the samples manufacturing. The stripes observed on films produced with mLLDPE A can be explained by slight differences on cooling air Polímeros, 25(6), 531-539, 2015

Table 5. Energy and output indexes. Samples mLLDPE A mLLDPE B Variation B/A (%)

Energy index (A.h/kg) 0.99 0.86 –13

Output index (bar.h/kg) 3.71 3.13 –16

Table 6. Optical properties. Samples mLLDPE A_35 mLLDPE B_35 mLLDPE A_60 mLLDPE B_60 mLLDPE A_100 mLLDPE B_100

Haze (%) 15.2 ± 2.5 10.2 ± 0.4 16.7 ± 2.5 9.8 ± 0.4 27.1 ± 1.0 11.2 ± 0.7

Gloss 45o 62.7 ± 5 60.4 ± 1 64.3 ± 4 66.8 ± 2 55.3 ± 4 69.6 ± 1

temperature. Film cooling is made with cold air inside and outside of the bubble. The air out of the bubble comes from small orifices, and, sometimes, temperature is not 100% homogeneous. Considering LLDPE polymers fast crystallization kinetics, these little differences are enough to produce discrepancies on film crystallization[29]. As the mLLDPE B sample possesses LCB, the crystallization kinetics is reduced, so the slight differences in cooling temperature does not disturb film homogeneity[19]. When the film thickness increases, haze values for the films produced with mLLDPE A also increase, while corresponding values for the films produced with mLLDPE B are at the same range. Gloss 45° values increase while the mLLDPE B-based films thickness increases. For the samples produced with mLLDPE A, the behavior is different. Gloss 45° values of the films with 35 µm and 60 µm are at the same range, but the film with 100 µm presents a smaller value 535

Dartora, P. C., Santana, R. M. C., & Moreira, A. C. F. for Gloss 45°. Usually it is desirable to produce films with low haze and high gloss for the packaging industry when the product does not need light protection[2,30]. The lower haze observed for mLLDPE B should be explained by the absence of stripes during blown film extrusion. Other factors that could affect haze are degree of crystallization and/or crystal structure and relaxation time of polymers. It is known that LCB increases relaxation time of

polymers[31,32]. Higher relaxation time allows the crystallization to occur under influence of stress elongation, causing it to form small, thin and oriented crystalline structures. Hence this film has lower haze and higher gloss[20]. The Elmendorf tear strength data are shown in Figure 5. Clearly, all samples exhibit higher values in the transversal direction (TD) when compared to the machine direction (MD). The morphological developments during blown

Figure 4. Blown film extrusion: (a) mLLDPE A; (b) mLLDPE B.

Figure 5. Elmendorf tear resistance: (a) mLLDPE A based films; (b) mLLDPE B based films. 536

Polímeros , 25(6), 531-539, 2015

The influence of long chain branches of LLDPE on processability and physical properties film process for LLDPE explain the trends of tear strength TD > tear strength MD[31]. Tear strength in the MD was extremely affected by film orientation[20]. TD and MD tear strengths can be compared using the ratio of the values obtained. Table 7 shows the aforementioned ratios. From the results above, one can deduce films produced with mLLDPE A have a better balance between MD and TD tear strength on all samples, showing a MD/TD ratio of approximately 1. On the other hand, mLLDPE B-based films have a sizeable difference between MD and TD tear strength and as a result their MD/TD ratios are smaller than 0.5 for samples with thicknesses of 35 µm and 60 µm and about 0.5 for the 100 µm-thick sample. It is known that films produced with LDPE resins have a higher tear resistance in the MD when film orientation is increased[20,31]. In the LDPE films, the twisted lamellae from adjacent row nuclei are strongly connected, which is responsible for the high MD tear. When the film orientation is increased, these connections get stronger, increasing even more the tear resistance. Films produced with LLDPE have the opposite Table 7. Tear strength MD/TD ratio. Thickness (μm) 35 60 100

mLLDPE A 0.85 0.82 0.88

mLLDPE B 0.31 0.32 0.51

behavior and consequently their tear strength is reduced when they are highly oriented because they usually have less oriented localized spherulite-like structures[20]. Based on this information, it can be said that the samples produced with the mLLDPE B resin present a higher level of orientation than the samples produced with mLLDPE A. When the film thickness increases, the difference between MD and TD tear strength is reduced. To increase film thickness, it is necessary to stretch it less during the manufacturing process, so the high orientation observed in the MD is reduced and, therefore, so are the differences between MD and TD tear resistances. In other words, film anisotropy is reduced. Figure 6b clearly shows tearing in the MD for mLLDPE B_35 sample. The sample mLLDPE A 35 exhibits a rough surface on tearing propagation in the MD. The same behavior was observed for the samples with thicknesses of 60 µm and 100 µm. The observed difference in tear resistance can be associated with the crystalline lamellar structure formed at film processing. As mentioned previously, LLDPEs in general have a less oriented localized spherulite-like structure but, when long chain branches are present, it is possible that the lamellae structure get more similar to the structure observed on LDPEs, which presents twisted lamellae from strongly connected adjacent row nuclei[20].

Figure 6. Elmendorf tear test specimens: (a) mLLDPE A_35; (b) mLLDPE B_35. Polímeros, 25(6), 531-539, 2015

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Dartora, P. C., Santana, R. M. C., & Moreira, A. C. F. Table 8. Films crystallinity measured with DSC analysis. Thickness (μm) 35 60 100

mLLDPE A 42.0 43.5 46.8

mLLDPE B 45.8 47.3 46.8

DSC results were used to calculate film crystallinity. The results are shown on Table 8. For sample mLLDPE A, it was observed that increasing film thickness, the crystallinity also increases. On the other hand, for sample mLLDPE B, the crystallinity degree was nearly the same. Film samples with higher MD stretch ratio (35 and 60 μm) presented higher crystallinity for sample mLLDPE B, then for sample mLLDPE A. The sample with lower MD stretch ratio showed the same crystallinity degree for both samples. These results showed that crystallinity degree is not the only explanation for optical properties differences obtained. The optical differences observed for the films with different thicknesses are probably caused by crystallites shape and size, according with previously publications[22,24,31,33,34].

4. Conclusions The presence of LCB was confirmed for one of the resins (mLLDPE B) evaluated in this study. The characterization of LCB in LLDPE was obtained by 13C NMR and rheology analyses. It was clearly shown specific chemical shifts at NMR to confirm the LCB presence in mLLDPE B, and also the possibility to quantify the level of LCB/10000C. Also, rheology studies results allowed the obtention of flow activation energy for the samples and the observation of a major difference in the values of flow activation energy. In agreement with literature, it was found that even small levels of LCB significantly altered the processability of the blown film resins. In particular, it was possible to measure the difference between resins though energy and output indexes. These data confirm that LLDPE with LCB are more prone to reduce costs during blown film extrusion. In relation to optical properties, it was possible to confirm the positive effect of LCB on haze and gloss, independently of the film thickness, contributing for better flexible packaging films. However, it was observed that the addition of LCB to LLDPE blown film resin resulted in a decrease in Elmendorf tear resistance. The presence of LCB is likely to produce higher levels of orientation on blown film and as a result of this a huge unbalance in tear resistance was observed. When the stretch ratio was reduced, samples presented better Elmendorf tear resistance, but a reduction in the stretch ratio results in lower productivity. The films here studied did not show significant crystallinity variations by changing film thickness. The introduction of LCB does not appear to provide an improvement for both processing and film performance, as it has been often suggested in the literature. Probably there is an optimum level of LCB that improves both processing and properties. Thence, we intend to further characterize both polymer and films samples to completely understand the LCB effect on crystallite formation during blown film process. 538

5. Acknowledgements The authors would like to thank the following for their significant contributions to this work: M. A. da Silva, B. E. S. Mendonça, F. P. dos Santos, S.S. Staub and C. Ellwanger. In addition, the authors would like to thank Braskem S.A. for support and permission to publish this work and CNPq for financial support.

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(LDPE). Pure and Applied Chemistry, 60(9), 1403-1415. http:// dx.doi.org/10.1351/pac198860091403. 24. Bueche, F. (1962). The physical properties of polymers. New York: John Wiley&Sons. 25. Bretas, R. E. S. & D’avila, M. A. (2000). Reologia de polímeros fundidos. São Carlos: UFSCar. 26. Dealy, J. M., & Wissbrun, K. F. (1990). Melt rheology and its role in plastics processing. London: Chapman&Hall. 27. Malkin, A. Y. (1994). Rheolgy fundamentals. Toronto: ChemTec. 28. Butler, T. I. (2005). Film extrusion manual: process, materials, properties. Norcross: Tappi Press. 29. Guerrini, L. M., Paulin, F. P. I., & Bretas, R. E. S. (2004). Correlação entre as propriedades reológicas, óticas e a morfologia de filmes soprados de LLDPE/LDPE. Polímeros: Ciência e Tecnologia, 14(1), 38-45. http://dx.doi.org/10.1590/ S0104-14282004000100012 30. Alvarez, V. B., & Pascall, M. A. (2011). Packaging. In J. W. Fuquay (Ed.), Encyclopedia of dairy sciences (pp. 16-23). New York: Academic Press. 31. Zhang, X. M., Elkoun, S., Ajji, A., & Huneault, M. A. (2004). Oriented structure and anisotropy properties of polymer blown films: HDPE, LLDPE and LDPE. Polymer, 45(1), 217-229. http://dx.doi.org/10.1016/j.polymer.2003.10.057. 32. Krishnaswamy, R. K., & Sukhadia, A. M. (2000). Orientation characteristics of LLDPE blown films and their implications on Elmendorf tear performance. Polymer, 41(26), 9205-9217. http://dx.doi.org/10.1016/S0032-3861(00)00136-1. 33. Dartora, P. C., Moreira, A. C. F., Stocker, M. K., & Santos, F. P. (2013). Análise de polietileno linear de baixa densidade metalocênico por microscopia: efeito de ramificações longas. In Anais do 12° Congresso Brasileiro de Polímeros. Florianópolis: Associação Brasileira de Polímeros. 34. Liu, Z. J., Ouyang, J., Zhou, W., & Wang, X. D. (2015). Numerical simulation of the polymer crystallization during cooling stage by using level set method. Computational Materials Science, 97, 245-253. http://dx.doi.org/10.1016/j. commatsci.2014.10.038. Received: Mar. 27, 2014 Revised: Mar. 26, 2015 Accepted: June 23, 2015

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http://dx.doi.org/10.1590/0104-1428.2152

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

Surface treated fly ash filled modified epoxy composites Uma Dharmalingam1, Meenakshi Dhanasekaran1, Kothandaraman Balasubramanian1 and Ravichandran Kandasamy1* 1 Polymer Chemistry Laboratory, Department of Rubber and Plastics Technology, Madras Institute of Technology Campus, Anna University, Chennai, Tamil Nadu, India

*ravi@mitindia.edu

Abstract Fly ash, an inorganic alumino silicate has been used as filler in epoxy matrix, but it reduces the mechanical properties due to its poor dispersion and interfacial bonding with the epoxy matrix. To improve its interfacial bonding with epoxy matrix, surface treatment of fly ash was done using surfactant sodium lauryl sulfate and silane coupling agent glycidoxy propyl trimethoxy silane. An attempt is also made to reduce the particle size of fly ash using high pressure pulverizer. To improve fly ash dispersion in epoxy matrix, the epoxy was modified by mixing with amine containing liquid silicone rubber (ACS). The effect of surface treated fly ash with varying filler loadings from 10 to 40% weight on the mechanical, morphological and thermal properties of modified epoxy composites was investigated. The surface treated fly ash was characterized by particle size analyzer and FTIR spectra. Morphological studies of surface treated fly ash filled modified epoxy composites indicate good dispersion of fillers in the modified epoxy matrix and improves its mechanical properties. Impact strength of the surface treated fly ash filled modified epoxy composites show more improvement than unmodified composites. Keywords: fly ash, poly dimethyl siloxane-co-amino propyl methyl siloxane, silane coupling agent, surfactant.

1. Introduction Organic-inorganic hybrid composites have gained immense attention from the researchers, as they combine the advantages of both inorganic solids such as high mechanical, thermal and structural stability and the characteristics of the organic molecules provide flexibility and functionality[1]. Many polymer composite industries are trying to use inorganic filler in polymer matrix to improve the mechanical properties and to reduce the cost of the final product. Fly ash, an absolutely low cost inorganic waste product obtained from thermal power stations, is available abundantly and poses environmental hazards and disposal problems[2]. Fly ash mainly comprises alumino silicate, calcium oxide and Ferric oxide, which is used in construction industry because of its advantages such as low density, low cost, strong filling ability, smooth spherical surface, small and well distributed internal stress in the products and good processability of the filled materials[3]. But the usage in polymer industry is limited due to its weak interfacial’s bonding between fly ash and polymer matrix and due to the low friction of the fly ash surface[4]. Surface treatment of fly ash with reactive silanes or surfactants, capable of graft formation is one of the principal methods for converting mineral of inorganic particulate filler into the materials bearing covalently bound functional groups and this treatment is used to enhance the mechanical properties of fly ash filled polymer composites. The surface area of fly ash filled polymer composite can be increased by reducing the particle size of fly ash and this enables uniform distribution in polymer matrix and improves the mechanical properties[5]. The surface treatment of fly ash with surfactant enhances the physical properties by avoiding particle-particle interaction and has a better distribution of fly ash within the polymer[6]. The use of silane coupling agent in polybutadiene rubber, improves the

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interaction between the rubber and fly ash and enhances the mechanical properties of the rubber[5]. Surface treatment of fly ash with acetone and silane improves impact properties than untreated fly ash in epoxy matrix[7]. Epoxy resin a thermosetting polymer has many advantages such as high strength and stiffness, good chemical and electrical resistances and low cost[8]. The use of unmodified epoxy resins based on bisphenol A-epichlorohydrin exhibits brittleness and low elongation after cure. To provide toughness enhancement, phase separation followed by the slow development of a two-phase morphology is critically important. Therefore, the modification of epoxy should be done to improve its toughening properties and also to get better the dispersion and interfacial bonding of fly ash in epoxy matrix[9]. In this research work modification of fly ash was done using i) surfactant sodium lauryl sulfate and ii) silane coupling agent glycidoxy propyl trimethoxy silane and by iii) size reduction of fly ash was done upto submicron level using high pressure pulverizer. The surface treated fly ash was dispersed in amine containing liquid silicone rubber modified epoxy matrix. The effect of surface treatment of fly ash with varying filler loadings from 10 to 40%weight on the mechanical, morphological and thermal properties of modified epoxy was investigated.

2. Materials and Methods 2.1 Materials Fly ash, Class F grade was procured from North Chennai Thermal Power Corporation (NCTPC) Chennai. Epoxy resin based on diglycidyl ethers of Bisphenol A (LY556) and hardener aliphatic amine triethyltetramine

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Surface treated fly ash filled modified epoxy composites (HY951) from Huntsman, surfactant sodium lauryl sulfate (SLS) from Merck, silane coupling agent Glycidoxypropyl trimethoxysilane (GPTS) from Sigma Aldrich and other chemicals were procured from the local source. Amine containing liquid silicone rubber (ACS) a poly [dimethyl siloxane-co-(3‑amino propyl) diethyl siloxane] having its grade AN102 was supplied by Resil Chemicals Ltd, Bangalore.

2.2 Modification of epoxy resin The amine containing liquid silicone rubber modified epoxy was prepared by mixing the epoxy resin with 3 grams of amine containing liquid silicone rubber and heated at 50 °C for 30 min with constant stirring to have the same concentration in all the parts. Then the mixture was allowed to cool at room temperature[9].

2.3 Surface treatment of fly ash i) The silane coupling agent Glycidoxypropyl trimethoxysilane (GPTS) (2 gram) was mixed with 100 ml of ethyl alcohol and 40 grams of fly ash was added. The mixture was stirred for the uniform distribution of the coupling agent to the filler. The mixing was done for about 45 min and it was filtered with the filter paper. The filtrate was then dried at 100 °C in a hot air oven for about 8 hrs[5]. The silane treated fly ash was designated as SIL. ii) Chemical modification of fly ash was done by taking 20 gram of fly ash mixed with 200 ml of distilled water. The surfactant sodium lauryl sulfate was added in the concentration of 2% wt and stirred for 6 hrs at temperature of 60 °C and then the samples were filtered and washed with distill water and dried in an oven for 48 hrs at 60 °C[10]. The sodium lauryl sulphate treated fly ash was designated as SLS. iii) The particle size of fly ash was reduced by using a high pressure pulveriser by maintaining 7 bar pressure and its particle size distribution was analysed. The size reduced fly ash was designated as SUM.

2.4 Preparation of composites The modified epoxy resin was mixed with unmodified fly ash in micron size (MIC), size modified fly ash in submicron level (SUM), sodium lauryl sulfate treated fly ash (SLS) and silane treated fly ash (SIL) at various loading (10, 20, 30 and 40%weight) were physically blended by stirring for 5-10 min. Then the hardener was added in the ratio of resin: hardener (10:1) and it was stirred for 10 min. to prepare a homogeneous blend. The mould set up was kept in room temperature for 24 hrs and then placed in a hot air oven heated at 100 °C for 3 hrs. The composition of epoxy with fly ash was shown in Table 1.

MB 3000 FT-IR spectrometer by making the samples as disc using potassium bromide. The morphology of fly ash and its composites were examined by means of scanning electron microscope (SEM-JEOL JSM 850) and the magnification was varied from 100 to 10000X. Energy dispersive spectroscopy (EDX) of fly ash was analyzed for the elemental analysis. Thermo gravimetric analyses were done using TG AnalyzerModel Q50, TA Instruments, from 30 to 600 °C under N2 atmosphere with a flow rate of 40-60 ml/min with heating rate of 20 °C/min. Mechanical properties such as flexural properties were carried out using UTM Shimadzu as per the ASTM D 790 standards at a test speed rate of 1 mm/min. Izod impact test was conducted on unnotched sample of size 60 × 11 × 3mm as per ASTM D-256 standard to evaluate the toughness of the composite. An impact load was applied through a 2-4 J machine pendulum and the amount of energy absorbed before fracturing the samples was determined.

3. Results and Discussion For good reinforcement the particle size should be finer and should have higher surface area.Using optical microscope, the particle size distribution of very fine particles of fly ash were determined. Here, required magnification of the particles is done, followed by selecting individual particles to determine their size. Then, in order to arrive at the average particle size a frequency table is drawn as shown in Table 2 and it was observed that 90% of particles were in the size ranging between 25-227 µm. The particle size distribution of pulverised fly ash is shown in Figure 1. The data was used to find out the mean particle size, which was found to be between 500-600 nm (ie) in sub micron size and was designated as SUM. The chemical composition of fly ash was determined by Energy dispersive X-ray microscope (EDX), which works on the principle that each element in the composition of fly ash emits peaks of varied frequencies. The chemical composition of fly ash used was given in Table 3. It reveals the abundant presence of silica as the main constituent along with alumina and traces of other oxides. The morphology of fly ash and surface treated fly ash was analyzed using scanning electron microscopy. Figure 2a shows the SEM images of fly ash particles which are in Table 1. Composition of epoxy and Fly ash composites. Composition 1 2 3 4 5

Epoxy resin

Fly ash

ACS

(g) 100 90 80 70 60

(g) 10 20 30 40

(g) 3 3 3 3 3

2.5 Characterization and testing The particle size distribution of the fly ash was measured by optical microscope using Biowizard Zoomaster IV optical microscope. The particle size distribution of pulverized fly ash particles was done using instrument Malvern Zetasizer Nano-series utilizing water as the dispersion medium. Fourier transform-infrared spectra of fly ash were recorded on ABB Polímeros, 25(6), 540-546, 2015

Table 2. Particle Size Distribution of Fly ash by Optical Microscope. codes 1 2 3

Range (micron) 25-227 227-420 428-630

% 90.75 8.9 1.02

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Dharmalingam, U., Dhanasekaran, M., Balasubramanian, K., & Kandasamy, R. spherical shape with agglomeration and irregularly shaped amorphous particles can be detected and it is due to sudden cooling from high temperature[11]. The SEM image of silane treated fly ash is shown in Figure 2b. It is observed that the fly ash particles have interaction with silane coupling agent and the particles are well distributed and has less agglomeration formation. Figure 2c shows the SEM images of sodium lauryl sulfate treated fly ash. It is observed that the sodium lauryl sulfate was well coated on the fly ash particle which was expected to have good interaction with polymer matrix and also less formation of agglomeration. Figure 3a, b shows the SEM images of plain epoxy and modified epoxy. It is observed that the fracture surface of plain epoxy shows a thin line indicating a brittle fracture. On addition of amine containing liquid silicone rubber, the Table 3. Chemical composition of Fly Ash (Class F). Compound SiO2 Al2O3 Fe2O3 CaO MgO

Concentration (%) 40 19 6 2 3

Figure 1. Particle size distribution of pulverised fly ash.

rubber particles initially dissolve and become dispersed at a molecular level in the epoxy and get precipitated when epoxy cross linking takes place. This gives the required two phase morphology with the formation of rubbery particles dispersed and bonded to the cross-linked epoxy matrix. The Figure 3c, d give the SEM images of fractured surfaces of epoxy fly ash (10% MIC) loaded with and without modification of epoxy, and from these images it can be observed that there is a more uniform dispersion of the fly ash in the modified epoxy composites than in the unmodified epoxy, though the particles seem smaller in the unmodified epoxy composite. It can be assumed that the modification of epoxy by silicone can lead to better dispersion and distribution of the fly ash particle. This has been discussed in a few publications (eg) Altaweel et al.[12] and Takahashi et al.[13]. Takahashi et al. observed that in silica filled epoxy composites, the filler distribution is better if the epoxy is modified by amine terminated silicone and this leads to an increase in flexural strength. It is to be noted that silicone is a rubbery polymer and a rubbery additive is expected to decrease flexural strength, but the flexural strength gets increased due to better polymer-filler interactions caused by the silicone modification of the epoxy. Altaweel et al.[12], while studying the effect of ACS on epoxy-fly ash composites, gives evidence for the above, through free volume measurements. The silane modified fly ash with modified epoxy composites seems to have a better dispersion of the fly ash as well as a lesser tendency to agglomerate, as shown in Figure 3e and the SLS modified fly ash with modified epoxy composites is shown in Figure 3f. Agglomeration tendency for fly ash particles seems to be present in SLS modified filler composite too (Figure 3e). The finest distribution of fly ash seems to be in the silane modified fly ash in modified epoxy in Figure 3f The FTIR spectra of unmodified fly ash is shown in Figure 4a.The peak at 3500-3000 cm–1 is assigned to O-H bonding, the Si-O-Si bond assymmetric stretching is observed at 1600 cm–1. The FTIR spectra of silane treated fly ash is shown in Figure 4b. The peaks at 3500-3000 cm–1 is assigned to O-H bonding and it is weak compared to pure fly ash due to the reaction between silane and fly ash particles. The O-H group in Si coupling agent condenses with O-H group of fly

Figure 2. (a) SEM of fly ash; (b) SEM of silane treated fly ash; (c) sodium lauryl sulphate treated fly ash. 542

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Surface treated fly ash filled modified epoxy composites

Figure 3. (a) SEM of unmodified epoxy; (b) SEM of modified epoxy; (c) SEM of epoxy fly ash (10%) micron without ACS modification; (d) SEM of epoxy fly ash (10%) micron with ACS modification; (e) SEM of silane modified fly ash with modified epoxy composites; (f) SEM of SLS modified fly ash with modified epoxy composites.

ash in an aqueous environment leading to the formation of covalent bond linkage for improving the resin wettability. The peaks at 2800 cm–1 are assigned to the C-H stretch in the propyl chain and the peak at1030 cm-1 is interpreted to the Si-O-Si bonds[14].The FTIR spectra of sodium lauryl sulphate treated fly ash particles are shown in Figure 4c. It is observed that the peaks at the range of 1000-700 cm–1 are assigned for the vibration of Si-O and Al-O bonds. There is a slight shifting of band in the range of 1000 cm-1 to the higher wavelength because of the interaction of polar active SLS with the Si-O and Al-O bonds in fly ash[6]. The FTIR spectra of plain epoxy is shown in Figure 5b. The peak observed at 3050 cm–1 is assigned to C-H stretching of the oxirane ring, the C-H stretching of aromatic and aliphatic group of methyl groups is observed at the range of 2960-2800 cm–1, the peak at 1600 cm–1 is assigned to stretching of C=C of aromatic ring, the stretching of C-C of aromatic ring is observed at 1509 cm–1, the peak at 1036 cm–1 is assigned to the stretching of C-O-C group of ether and the stretching of C-O group of oxirane ring was observed at 900 cm–1[15].The FTIR spectra of the amine containing liquid silicone rubber as modifier for epoxy is shown in Figure 5c. The stretching vibration of Si-O bond of Si-OH is observed at 768 cm–1 and the stretching vibration of Si-CH3 is observed around 1200 cm–1.The stretching vibration of Si-O is observed at 1010 cm–1 and the peak at Polímeros, 25(6), 540-546, 2015

Figure 4. FTIR of (a) Fly ash; (b) Silane modified Fly ash; (c) SLS modified Fly ash.

2950 cm–1 is assingned to the stretching vibration of C-H group[16]. The FTIR spectra of amine containing liquid silicone rubber modified epoxy (ACS 3phr) is shown in Figure 5a. It is observed that a low intense peak is found in the range of 3500 cm–1, because 1:8 part of amine group is present in the modifier and this amine group reacts with the 543

Dharmalingam, U., Dhanasekaran, M., Balasubramanian, K., & Kandasamy, R.

Figure 6. Structure of Liquid Silicone Rubber.

Figure 5. FTIR of (a) ACS modified epoxy; (b) pure epoxy; (c) Amine containing Silicone modifier.

epoxide group and some O-H groups are formed. Then the modification at the peak of 900 cm–1 is due to C-O group of the oxirane ring opened by amine group. The structure of amine containing liquid silicone rubber is shown in Figure 6. The flexural strength of untreated and surface treated fly ash filled with modified epoxy composites is shown in Figure 7. It is observed that on increasing the fly ash loading from 10 to 40% weight, the flexural strength is found to be decreasing. Silane treated fly ash composites show an improvement in flexural strength compared to SLS treated fly ash and size reduced fly ash (500 to 600 nm) modified epoxy composites. Altweel et al.[12] too reported that fly ash content beyond 10% reduced flexural strengths in such composites. Figure 8 shows the flexural modulus of untreated and surface treated fly ash filled with modified epoxy composites. The flexural modulus is found to be higher at 10%weight loading of fly ash for both untreated and treated fly ash. There is an increase in flexural modulus for 10% fly ash filler loading; comparatively treated fly ash has a higher flexural modulus than the untreated fly ash loading. The decrease in the flexural modulus of the composites after 10% filler is due to the poor interaction between the fly ash and epoxy matrix[10]. It may be due to the agglomeration of the fillers creating filler-filler interaction than polymer filler interaction. The unnotched Izod impact strength of epoxy fly ash composites is shown in Figure 9. It is observed that the impact strength increases up to fly ash loading 30%weight and at 40% weight loading there is a decrease in the impact strength because of the high filler content, the polymer availability for filler interaction may be low[9]. Comparatively, surface modified fly ash composites show higher impact strength than the untreated fly ash because of the better dispersion. The thermal stability and thermal degradation patterns of the composites are studied by thermo gravimetric analysis. The thermograms of modified epoxy and epoxy – fly ash composites are shown in Figure 10a, b. From the figure it is observed that the degradation for resins occurrs around 544

Figure 7. Flexural strength of surafce treated fly ash filled modified epoxy composites.

Figure 8. Flexural modulus of surafce treated fly ash filled modified epoxy composites.

250 °C, with 5-8% weight loss. This fact is attributed to the presence of impurities. The onset of degradation temperature starts at the temperature above 300 oC for almost all the composites. The char residue is high for 40% weight fly ash loading than 10% weight fly ash content. Epoxy has minimum char residue around 5% weight because of the presence of silicone in modified epoxy. It can be confirmed that the Polímeros , 25(6), 540-546, 2015

Surface treated fly ash filled modified epoxy composites stability of the blends is slightly increased with increasing fly ash content due to the high siliceous content of fly ash.

4. Conclusions

Figure 9. Impact strength of surafce treated fly ash filled modified epoxy composites.

Modification of plain epoxy is done by adding amine containing liquid silicone rubber with the aim of reducing fly ash agglomeration formation and improving the dispersion of fly ash in modified epoxy matrix. Comparative studies were also made by surface treatment of fly ash using i) sodium lauryl sulphate, ii) GPTS and iii) size reduction by high pressure pulverizer. SEM images show that silane coupling formation compared agent treated fly ash particles are well distributed and has less agglomeration to SLS treated fly ash. Silane treated fly ash composites show an improvement in flexural strength compared with SLS treated fly ash and size reduced fly ash filled epoxy composites (upto 10% fly ash loading). It is also observed that in SEM images of the composites and impact strength of fly ash composites increase upto 30% fly ash loading and on further loading it decreases. Silane treated fly ash composites show an increase in impact strength than MIC, SUM and SLS treated fly ash composites. The thermal stability of the epoxy fly ash composites is increased with the addition of fly ash. Based on the above observations, it may be stated that the addition of silane treated fly ash at 10% by weight of the resin, can be considered as a good composition for achieving good mechanical properties in modified epoxy composites.

5. Acknowledgements The authors gratefully acknowledge Anna University for financial support to carry out this research work under Anna Centenary Research fellowship.

6. References

Figure 10. (a) TGA graph of ACS modified epoxy fly ash composites (10%) loading; (b) TGA graph of ACS modified epoxy fly ash composites (40%) loading Polímeros, 25(6), 540-546, 2015

1. Akinci, A. (2009). Mechanical and morphological properties of basalt filled polymer matrix composites. Archives of Materials Science and Engineering, 35, 29-32. 2. Devi, M. S., Murugesan, V., Rengaraj, K., & Anand, P. (1998). Utilisation of fly ash as filler for unsaturated polyester resin. Journal of Applied Polymer Science, 39(7), 1385-1391. http:// dx.doi.org/10.1002/(SICI)1097-4628(19980815)69:7<1385::AIDAPP13>3.0.CO;2-T. 3. Sreekanth, M. S., Joseph, S., Mhaske, S. T., & Mahanwar, P. (2009). Effect of particle size and concentration of fly ash on properties of polyester thermoplastic elastomer composites. Journal of Minerals and Materials Characterization and Engineering, 8(3), 237-248. 4. Yang, Y. F., Gai, G. S., Cai, Z. F., & Chen, Q. R. (2006). Purification and surface modification of fly ash. Journal of Hazardous Materials B, 133, 276-282. 5. Alkadesi, N. A. N., Hundiwale, D. G., & Kapadi, U. R. (2004). Effect of coupling agent on the mechanical properties of fly ash–filled polybutadiene rubber. Journal of Applied Polymer Science, 91(2), 1324-1328. http://dx.doi.org/10.1002/app.13280. 6. Nath, D. C. D., Bandyopadhyay, S., Gupta, S., Yu, A., Blackburn, D., & White, C. (2010). Surface-coated fly ash used as filler in biodegradable poly(vinyl alcohol) composite films: part 1-The modification process. Applied Surface Science, 256(9), 2759-2763. http://dx.doi.org/10.1016/j.apsusc.2009.11.024. 545

Dharmalingam, U., Dhanasekaran, M., Balasubramanian, K., & Kandasamy, R. 7. Kulkarni, S. M., & Kishore, (2002). Effects of surface treatments and size of fly ash particles on the compressive properties of epoxy based particulate composites. Journal of Materials Science, 37(20), 4321-4326. http://dx.doi.org/10.1023/A:1020648418233. 8. Singla, M., & Chawla, V. (2010). Mechanical properties of epoxy resin – fly ash composite. Journal of Minerals & Materials Characterization & Engineering, 9(3), 199-210. http://dx.doi.org/10.4236/jmmce.2010.93017. 9. Mohammed Altaweel, A. M. A., Ranganathaiah, C., Kothandaraman, B., Raj, J. M., & Chandrashekara, M. N. (2011). Characterization of ACS modified epoxy resin composites with fly ash and cenospheres as fillers, mechanical and microstructural properties. Polymer Composites, 32(1), 139-146. http://dx.doi.org/10.1002/pc.21030. 10. van der Merwe, E. M., Prisloo, L. C., Kuger, R. A., & Mathebula, L. C. (2011). Characterization of coal fly ash modified by sodium lauryl sulfate. 2011 World of Coal Ash Conference (pp. 01-17). Denver: University of Kentucky. Retrieved in 10 May 2011, from http://www.flyash.info/2011/083-vanderMerwe-2011. pdf 11. Satheesh Rajaa, R., Manisekar, K., & Manikandan, V. (2013). Effect of fly ash filler size on mechanical properties of polymer matrix composites. International Journal of Mining, Metallurgy &. International Journal of Mining, Metallurgy & Mechanical Engineering, 1(1), 34-37. Retrieved in 05 January 2013, from http://www.isaet.org/images/extraimages/IJMMME%20 0101008.pdf 12. Altaweel, A. M. A. M., Ranganathaiah, C., & Kothandaraman, B. (2009). Mechanical properties of modified epoxies as related

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to free volume parameters. The Journal of Adhesion, 85(4-5), 200-215. http://dx.doi.org/10.1080/00218460902881774. 13. Takahashi, T., Nakajima, N., & Saito, N. (1989). Morphology and mechanical properties of rubber-modified epoxy systems. Amino-terminated polysiloxane. In C. K. Riew (Ed.), Rubber toughened plastics (Vol. 222, pp. 243-261). Washington: American Chemical Society. 14. Wu, G., Gu, J., & Zhao, X. (2007). Preparation and dynamic mechanical properties of polyurethane-modified epoxy composites filled with functionalized fly ash particulates. Journal of Applied Polymer Science, 105(3), 1118-1126. http:// dx.doi.org/10.1002/app.26146. 15. González González, M., Cabanelas, J. C., & Baselga, J. (2012). Applications of FTIR on epoxy resins - identification, monitoring the curing process, phase separation and water uptake. In T. Theophile (Ed.), Infrared spectroscopy-materials science engineering and technology (pp. 261-284). Rijeka: InTech Publication. 16. Launer, P. J. (1987). Infrared analysis of organosilicon compounds: spectra-structure correlations. In B. Arkles (Ed.), Silicon compounds: silanes & silicones (pp. 100-103). Morrisville: Gelest. Received: Mar. 19, 2015 Accepted: July 01, 2015

Polímeros , 25(6), 540-546, 2015

http://dx.doi.org/10.1590/0104-1428.2097

Synthesis and characterization of pH and temperature responsive poly(2-hydroxyethyl methacrylate-co-acrylamide) hydrogels Manuel Rapado1 and Carlos Peniche2* Center of Technological Applications and Nuclear Development – CEADEN, Havana, Cuba 2 Center of Biomaterials, University of Havana, Havana, Cuba

*peniche@biomat.uh.cu

Abstract Acrylamide/2-hydroxyethyl methacrylate hydrogels were prepared by simultaneous radiation-induced cross-linking copolymerization of acrylamide (AAm), 2-hydroxyethyl methacrylate (HEMA) and water mixtures at a radiation dose of 10 kGy. Hydrogels were characterized by infrared spectroscopy. Dynamic and equilibrium swelling of hydrogels in water and in buffer solutions were investigated. They were sensitive to pH and temperature. Swelling was non-Fickean and increased with increasing the acrylamide content. Temperature dependence of the equilibrium water uptake of copolymers exhibited a discontinuity around 35 °C resulting from the weakening of the hydrogen bonds between the hydroxyl groups of HEMA and the amide groups of AAm. The thermodynamic and network parameters derived from swelling and mechanical measurements are compared and discussed. They exhibit a strong dependence on the AAm content in the hydrogel. These hydrogels can be considered for applications in fields requiring environmentally responsive hydrogels such as medicine, pharmacy and bioengeneering. Keywords: acrylamide, hydrogel, hydroxyethyl methacrylate, radiation copolymerization, swelling kinetics.

1. Introduction Hydrogels are three dimensional hydrophilic polymer networks that are capable to absorb considerable amounts of water, but do not dissolve when brought into contact with aqueous solutions[1]. In water, hydrogels swell to equilibrium while preserving their shape. The ability to swell and the extent of swelling of hydrogels are governed by two factors: the hydrophilicity of the polymer and the cross-link density of polymer chains[2]. Some hydrogels swell and contract in response to external stimulus such as pH, temperature and ionic concentration. These “smart” hydrogels have found a great variety of applications in biomedicine, pharmacy, biotechnology, bioengineering and agriculture, among others. Hydrogels based on 2-hydroxyethyl methacrylate (HEMA) are usually prepared by bulk polymerisation or polymerisation in solution in the presence of free radical initiators and cross-linking agents[3,4]. They have been used in controlled release drug delivery systems, contact lenses, artificial implants, burn dressings, among others due to their biocompatibility, hydrophilicity, softness, and permeability[3]. However, in some cases applications are limited due to poor mechanical properties or by low release capacity. Many approaches have been developed to improve their mechanical strength. These include using special co-monomers, changing the type and concentration of the cross-linking agent, optimizing polymerization conditions, forming interpenetrating polymer networks, and radiation cross-linking, among others[5-7]. The main advantage of radiation over conventional chemical methods for the preparation of biomaterials relies in its ability to modify macromolecules under fairly mild conditions, without additives. Besides,

Polímeros, 25(6), 547-555, 2015

radiation solves the need of sterilization providing a sterile product without toxic initiator residues[8,9]. Maximum swelling in water of pHEMA is thermodynamically limited to about 40%[10]. Therefore, in order to increase its swelling degree HEMA has been copolymerized with very hydrophilic monomers, such as methacrylic acid (MAAc) [11] and acrylamide (AAm). Copolymers of HEMA and acrylamide (AAm) have been prepared and studied by various authors[11-14]. They have shown that there is a significant dependence of swelling on the copolymer composition and that the copolymers prepared are pH responsive. The mechanical properties of poly(HEMA-co-AAm) hydrogels have been reported before[11,15]. However, the results are given for only one HEMA/AAm composition and the copolymers were prepared in presence of a chemical cross-linker. In order to evaluate the effect of composition on the properties of this copolymer system, in this work we prepare poly(HEMA-co-AAm) hydrogels with different compositions by γ-irradiation and characterize them in terms of morphology, swelling behaviour and mechanical properties. The thermodynamic and network parameters derived from swelling and mechanical measurements are compared and discussed.

2. Materials and Methods 2.1 Materials Acrylamide (AAm), 98.5% pure, was obtained from BDH (England) and used as received. 2-hydroxyethyl methacrylate (HEMA), 99% pure was purchased from Fluka (Switzerland).

547

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

Rapado, M., & Peniche, C. HEMA was purified by passing through a column filled with a HQ/MEHQ resin from Polysciences (Eppelheim, Germany) for inhibitor removing, and immediately before use it was purified by vacuum distillation (at 10 Pa) at 50 °C. Only the middle fraction collected was used.

Spectrophotometer (Nicolet Instrument Corp., Madison, WI). Spectra were obtained with a resolution of 2 cm–1 and were averaged over 100 scans. Samples (5 mg) were thoroughly dried and ground with KBr. Discs were prepared by compression under vacuum.

2.2 Copolymer synthesis

2.6 Morphologycal analysis

The poly(HEMA-co-AAm) hydrogels were prepared by simultaneous radiation–induced copolymerization and self‑bridging of HEMA and AAm in aqueous solution without the use of any cross-linker. The reactants were dissolved in double distilled water. The monomer feed compositions for the hydrogels syntheses are listed in Table 1. The monomer solutions were sealed in 10 mm diameter and 100 mm long glass bulbs under N2O- saturated atmosphere. They were irradiated with a 10 kGy dose at 25 °C. The dosimetry control for each exposition was determined using Fricke and ceric‑cerous dosimetry. Unreacted monomer and uncross‑linked polymers were removed by washing the gels for two weeks in distilled water. They were dried to constant weight in vacuum oven at 40 °C.

The morphologies of non-lyophilized and lyophilized hydrogels were determined using an environmental scanning electron microscope (ESEM) FEI Model Quanta 200 Philips/FEI (Redwood City, CA).

2.3 Hydrogel composition The chemical composition of hydrogels was calculated from C, H, and N elemental microanalysis performed on a CHNS-OEA1108-Elemental Analyzer, from Carlo Erba Instruments (Norwich, UK).

2.4 Density measurements The densities of non-lyophilized hydrogels were determined with a pycnometer using cyclohexane as non‑solvent. All density measurements were performed at 25 °C. They were calculated using the following equation[16]:

mxg ρcyc (1) ρ xg = m1 + mxg − m2

wheremxg is the mass of hydrogel, m1 the mass of pycnometer with cycloexane, m2 the mass of pycnometer with cycloexane and hydrogel and ρcyc the density of cycloexane (0.78 g∙cm−3). Experiments were carried out in triplicate and the mean values for the densities are listed in Table 1.

2.5 FTIR analysis FTIR spectra were obtained with KBr discs and recording the spectral range from 4000 to 500 cm–1 by using a Nicolet AVATAR 330 Fourier-transformed Infrared

2.7 Swelling experiments Dynamic swelling experiments were carried out by placing previously dried hydrogel discs of about 0.7 cm diameter and 0.2 cm thick in vials with 50 mL of double distilled water at 25 °C or in buffer solutions of various pHs (pH 2, 4, 6, 7.4, 8.5, 10, and 12) at 37 °C. The water uptake, W, was calculated by measuring the weight gain of the sample at different times after carefully wiping the surface with a filter paper. It was reported as W =

mt − m0 mt = −1 (2) m0 m0

where mt is the weight of the swollen disc at time t and m0 is the weight of the dry disc. All swelling experiments were performed in triplicate. The water penetration velocity (v) into hydrogels was determined by the weight gain method described by Davidson and Peppas[17]. To calculate v, Equation 3 was used:  1   dWg  = v   ×  2 A   dt 

(3)

where (dWg/dt) is the initial slope of the water uptake versus time curve, ρ denotes is the density of water at the measured temperature, A is the area of one face of the disc, and the factor 2 accounts for the fact that diffusion takes place through both faces.

2.8 Preparation of buffer solutions Buffers at pH 2, 4, 6, 7.4, 8.5, 10 and 12 were prepared following the procedure described elsewhere[18]. Briefly, 2.4732 g of boric acid were added to a mixture of 2.3 mL of glacial acetic acid and 2.7 mL of phosphoric acid and the resulting solution was diluted with double-distilled water to

Table 1.Composition of the reaction mixtures and gel fraction (%) for the radiation copolymerization of HEMA and AAm. Reaction mixture composition Sample polyHEMA M1 M2 M3 M4 M5 polyAAm

HEMA

(wt-%) AAm

H2O

40 40 40 23 10 5 0

0 5 12 18 10 24 24

60 55 48 59 80 71 76

molar fraction Feed Copolymer* -0.83 0.63 0.41 0.36 0.10 --

-0.81 0.65 0.41 0.37 0.11 --

Gel fraction (%)

Polymer density (g/cm3)

97 ± 0.4 99 ± 0.2 97 ± 0.3 99 ± 0.6 97 ± 0.8 98 ± 0.5 95 ± 0.3

1.31 ± 0.02 1.36 ± 0.02 1.38 ± 0.04 1.34 ± 0.05 1.27 ± 0.01 1.32 ± 0.02 1.30 ± 0.05

*Composition determined by elemental microanalysis using Equation 8.

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Polímeros , 25(6), 547-555, 2015

Synthesis and characterization of pH and temperature responsive poly(2-hydroxyethyl methacrylate-co-acrylamide) hydrogels a total volume of 1.0 L. Afterwards 50 mL portions of this solution were taken, and the pH was adjusted to the desired pH by adding an appropriate amount of 0.2 M NaOH.

2.9 Mechanical properties Measurements were performed by imposing a strain on the sample by uniaxial compression and measuring the resulting stress using a texture analyser TA-HD (Stable Micro Systems, UK). The apparatus was equipped with an 18 mm diameter cylindrical probe running at 0.1 mm∙s–1. All mechanical measurements were conducted in a constant temperature room at 25 ± 0.1 °C. Hydrogel samples of 18 mm diameter and 5-6 mm thick were exposed to unidirectional strain. To this end the gel discs were placed on the plate of the compression cell after gently wiping their surface with filter paper. Strain was applied up to 40% compression. Measurements were repeated five times on different discs prepared under identical conditions. The stress (σ) and strain (ε) were obtained according to Equations 4 and 5. P σ = (4) A l ε = F (5) l0

Where P represents the compression force, A the area of the cylinder and l0 and lF are the initial and final length, respectively. Elasticity modulus (E) values of hydrogels were calculated directly from the initial slope of the stress‑strain curves. The average molar mass between cross-links was estimated from the elasticity modulus by the following the equation[16]: = Mc

3ρ p RT E

(6) ϕ1/3 2m

where R is the gas constant, T the absolute temperature, ρp the polymer density, and φ2M the equilibrium volume fraction of polymer in the swollen state. The volume fraction of the swollen polymer was calculated using the following relation:  ρp  m  ρp  ϕm = 1 +  s  −   ρs  m0  ρs 

−1

(7)

In which ρpand ρs are are the densities of polymer and solvent, respectively, m0 and ms are the mass of polymer before and after swelling, respectively. Determinations were performed at 25°C.

3. Results and Discussions 3.1 Copolymer composition When HEMA/AAm/H2O mixtures are γ-irradiated free radicals are generated in the aqueous media. Random reactions of these radicals with the monomers lead to the simultaneous polymerization and cross-linking of monomers giving rise to cross-linked poly(HEMA-co-AAm) hydrogels. In this study a series of copolymerization reactions were carried out in water at different monomer feed compositions, while keeping the overall comonomer concentration between 2 and 5 M. For the preparation of mechanically stable hydrogels, the ternary mixtures of HEMA/AAm/H2O were irradiated at a radiation absorbed dose of 10 kGy with a Co-60 gamma source isotope at 25±1 °C. The copolymerization reaction can be represented, in a simplified form, according to the following equation (Figure 1): A series of poly(HEMA-co-AAm) hydrogels were obtained by varying the monomer feed weight ratio (HEMA/AAm) from 40:5 to 5:24. The considerations for selecting the feed compositions used in the present work were the solubility of AAm in HEMA and the shape stability of obtained hydrogels in their fully swollen state. The composition of the network chains in the hydrogels was estimated from elemental analysis. The copolymer composition can be represented as (C3H5NO)X(C6H10O3)1–X, and it can be readily shown that the molar fraction x of AAm in the copolymer can be calculated using the following relation: C (%)

N (%)

( 36 − 18 x ) (8) (7x)

Where C (%) and N (%) are the corresponding carbon and nitrogen contents in the copolymer as determined by elemental microanalysis. The results are shown in Table 1. The copolymerization parameters of HEMA (1) with AAm (2) are r1 = 1.84 and r2 = 0.41[19], but as it can be seen in Table 1, due to the high degrees of conversion attained in the present work, in all cases the copolymer composition is almost the same as that of the feed. As it can be seen in the table, there are no important differences in the samples densities with composition.

3.2 Spectral characterization The infrared spectra of polyHEMA, polyAAm and poly(HEMA-co-AAm) hydrogels M1 and M3 obtained from the corresponding monomer solutions (Table 1) at

Figure 1. Schematic representation of the radiation copolymerization reaction of HEMA and AAm. Polímeros, 25(6), 547-555, 2015

549

Rapado, M., & Peniche, C. constant dose of 10 kGy of gamma radiation are shown in Figure 2. Copolymers with other HEMA/AAm compositions exhibited similar qualitative spectra and are not shown. PolyAAm presents a distinctive absorption band at 1657 cm–1 (C=O stretching, amide I). The N–H stretching absorbs in the range 3500–3100 cm–1 and the band at 1460 cm−1 was associated with the stretching vibration of acyl amino groups[20]. On the other hand, polyHEMA exhibits a sharp absorption band at 1730 cm–1, which is characteristic of the C=O stretching. The absorption band at 2950 cm–1 with a weak shoulder is attributed to C-H stretching vibration of methylene groups. The wide O-H stretching absorption band appearing at 3500–3220 cm–1 indicates the presence of hydrogen bonding between hydroxyl groups of polyHEMA[14]. The IR spectrum of M1 and M3 copolymers exhibit all the characteristic bands of the corresponding homopolymers. It is worth mentioning the appearance of a broad absorption band peaking at 3220 cm–1 in the spectrum of the copolymers indicating the formation of hydrogen bonding, most probably between the –OH groups of HEMA and the –NH2 of the amide groups of AAm. The absorption bands observed at 1277 and 1460 cm–1, characteristic of the methyl and amide groups of HEMA and AAm, respectively, are also present in the spectra of M1 and M3[14]. It can be appreciated that the relative intensity of the main polyAAm absorption bands are greater in the acrylamide richer copolymer M3.

Figure 2. FTIR spectra of polyHEMA, polyAAm and poly(HEMA‑co-AAm) hydrogels M1 and M3.

3.3 Swelling of hydrogels A very important intrinsic characteristic property of hydrogels is their ability to swell when immersed in water or aqueous solutions[21].In the present study the effect of copolymer composition, pH and temperature on swelling for HEMA/AAm copolymers obtained by irradiation at 10 kGy were investigated. 3.3.1 Swelling as function of copolymer composition Swelling of polyHEMA, polyAAm and poly(HEMA‑co‑AAm) hydrogels of various compositions (Table 1) was followed at 25 °C in water for a long time to ensure equilibrium. The resulting swelling curves are shown in Figure 3 and the equilibrium water uptakes for all compositions are listed in Table 2, where a strong dependence of the swelling degree on composition is observed. As expected, polyAAm experienced higher maximum swelling (21.25 g/g) than polyHEMA (1.023 g/g), due to the greater hydrophilicity of the former. It has been pointed out that the low swelling of polyHEMA is due to intra- and intermolecular hydrogen bonding. In HEMA-rich copolymers swelling is controlled by the intermolecular hydrogen bonding between hydroxyl and amide groups and intramolecular hydrogen bonding between amide groups[12]. As the copolymers became richer in AAm, equilibrium swelling increased from 1.444 g/g for M1

Figure 3. Swelling at 25 °C of polyHEMA, polyAAm and poly(HEMA-co-AAm) hydrogels discs of different compositions (Table 1).

Table 2. Equilibrium water uptake and diffusion parameters of polyHEMA, polyAAm and poly(HEMA-co-AAm) hydrogels discs of different compositions (see Table 1). Polymer polyHEMA M1 M2 M3 M4 M5 polyAAm

550

W∞

Penetration velocity

(gH2O/g polymer)

(×10 cm/s)

1.023 1.444 2.364 4.504 9.776 20.62 21.25

3.65 ± 0.28 1.07 ± 0.25 2.93 ± 0.55 3.75 ± 0.27 3.90 ± 0.16 5.55 ± 0.38 9.57 ± 0.22

-5

Diffusion parameters (Equation 14) n 0.55 0.48 0.74 0.56 0.59 0.66 0.66

(k × 102) 2.9 2.4 0.44 1.7 1.2 0.77 0.20

Da (cm2/s) 2.188 × 10–6 1.241 × 10–7 1.630 × 10–7 3.451 × 10–7 4.159 × 10–7 4.651 × 10–7 5.312 × 10–6

Polímeros , 25(6), 547-555, 2015

Synthesis and characterization of pH and temperature responsive poly(2-hydroxyethyl methacrylate-co-acrylamide) hydrogels (HEMA/AAm 40:5) to 20.62 for M5 (HEMA/AAm 5:24). It is interesting to note in Table 2 that the initial solvent penetration velocity is higher in polyHEMA than in the copolymers with low AAm content (M1 and M2). This indicates that hydrogels of these two copolymers have a more compact structure than polyHEMA, which hinders the penetration of water during the initial swelling stage. Swelling kinetics of HEMA/AAm hydrogels obtained by γ-irradiation was evaluated using the Ritger–Peppas equation[22] (Equation 9) in order to describe the mechanism of water diffusion in the gels. Wt = kt n (9) W∞

where Mt/M∞ is the fractional water uptake, k is a kinetic constant, t is the swelling time and n is the diffusional exponent that can be related to the solvent transport mechanism. For a cylindrical sample, when n = 0.45, the water transport mechanism is Fickian diffusion. When n = 1, Case II transport occurs, leading to zero-order release. When the value of n is between 0.45 and 1, anomalous (non-Fickian) transport is observed. This mathematical model is valid only for the first 60% of the total solvent uptake. The values of the diffusional exponents, n and diffusion constants, k, can be calculated from slope and intercept of the logarithmic plot of the data in Figure 4 according to:

are shown in Figure 4. The water uptake of polyHEMA was small and almost pH independent in the whole pH interval studied, whereas the swelling degree of polyAAm was much higher and increased with increasing pH. This increase became more prominent as the pH raised from 6 to 12. This is the result of the basic hydrolysis of the amide groups of acrylamide, with formation of carboxylate ions. These negatively charged carboxylate ions are highly solvated and the repulsion between them provokes the expansion of the hydrogel network with the consequent increase in water uptake. As expected, the swelling capacity of the copolymer hydrogels increased with increasing their AAm content and they all exhibited the same basic hydrolysis effect on swelling at pH higher than 6, which became more significant as the AAm content in the hydrogel increased. A similar hydrolysis effect on swelling has been reported for other acrylamide based hydrogels[24]. An interesting feature of the system studied that has not been previously reported is the temperature dependence of W∞, the equilibrium water uptake of poly(HEMA-co-AAm) hydrogels. A plot of lnW∞ against the reciprocal of the swelling temperature is shown in Figure 5. Two straight lines with different slopes are obtained for all hydrogels in the

Wt ln = lnk + nlnt (10) W∞

They are reported in Table 2. In all cases the diffusional exponent n > 0.45, is indicating that the transport mechanism is non-Fickean. This anomalous behaviour was also obtained for HEMA/AAm hydrogels prepared by redox polymerization employing N,N’-mehthylenebisacrylamide as cross-linker. It is generally attributed to a slow relaxation rate of the polymer matrix[12]. The estimation of diffusion coefficients for non Fickiansorption processes is not straightforward. However, reasonably values of diffusion coefficients can be obtained by using the following simple equation[23]: (11)

Figure 4. Equilibrium water content of polyHEMA, polyAAm and M1, M3, M4 and M5 hydrogels at different pH values at constant ionic strength (I = 0.72 M) and 37 °C.

where Da is the apparent average diffusion coefficient in cm2∙s–1, l is one half of the cylindrical sample diameter in cm, t is the time at which the swelling is one half the equilibrium value (Mt/M∞= 0.5). The Da values reported in Table 2 are approximately one order of magnitude smaller for the copolymers as compared to the respective homopolymers. Probably the intramolecular bonding between the hydroxyl groups of HEMA and the amide groups of AAm in the copolymers are responsible for this decrease in Da. However, the general trend of the apparent average diffusion coefficients is to increase as the AAm content in the copolymer hydrogels increases. The pH dependence of the swelling capacity of polyHEMA, polyAAm, M1, M3, M4 and M5 hydrogels obtained with an irradiation dose of 10 kGy was evaluated in a wide pH interval at constant ionic strength (I = 0.72 M) and 37 °C. The equilibrium water uptakes for each composition and pH

Figure 5. Temperature dependence of the equilibrium water content, W∞, for poly HEMA, polyAAm and M1, M4 and M5 copolymers.

Da =

0.049

(t / 4l 2 )1/ 2

Polímeros, 25(6), 547-555, 2015

551

Rapado, M., & Peniche, C. temperature range studied. This behavior can be analyzed taking into consideration the Gibbs-Helmholtz equation for the infinite or equilibrium water content[25]: dln (W∞ ) d (1/ T )

= −

∆H m (12) R

where R is the gas constant and ΔHm is the enthalpy of mixing between the dry polymer and an infinite amount of water. The positive slope of the straight lines observed in Figure 5 for polyHEMA and M1 indicates an exothermic mixing process. Using Equation 12 the values of ΔHm were obtained. They are listed in Table 3 where it can be appreciated that the absolute values of ΔHm in the low temperature range (T < 35°C) are quite small. The small negative values of the enthalpy of mixing obtained for polyHEMA and M1 have been reported before for this homopolymer[26] and its copolymers with hydrophobic monomers[26,27]. The sudden change in the slope of the straight lines is interpreted in terms of the weakening of the hydrogen bonds between the OH the hydroxyl groups of HEMA and the amide groups of AAm in the copolymers. These hydrogen bonds act as physical cross-links that strengthen the structure and oppose swelling. Apparently above 35 °C the thermal energy is sufficiently high to lose these bonds allowing increasing the water uptake capacity of acrylamide and acrylamide richer copolymer hydrogels. A different swelling‑temperature behaviour of HEMA/AAm hydrogels prepared by γ-irradiation at 50 kGy in the presence of trithioglycolic acid (TTGA) was reported by El-Din and El-Naggar[14]. In contrast with the present findings their hydrogels shrank by heating above 25-30 °C. They concluded that these systems posses a lower critical solution temperature (LCST) at about 25 °C as a result of the formation of hydrophobic interchain bonding by copolymerization between TTGA and AM or HEMA.

3.4 Solubility parameters of HEMA/AAm copolymer hydrogels

Swelling experiments allow calculating the molecular weight between cross-links, Mc, which is the basic parameter that characterizes the structure of a hydrogel network. This parameter describes the average molecular weight of polymer chains between two consecutive junctions. These junctions include chemical cross-links, physical entanglements, crystalline regions, and polymer complexes (hydrogen bonding, salt bonds). Mc strongly influences the physical and mechanical properties of the hydrogel. Therefore, its determination is of great practical significance. It can be calculated according to the Flory-Rehner equation for the swelling of a perfect network[28,29]: M c = −ρ pV1

1/3 (1−2 / f ) ϕ2/3 2 r ϕ2 m (14) ln (1− ϕ2m ) + ϕ2m + χϕ22m

In Equation 14 ρP is the density of the polymer (in g∙cm–3), V1 is the molar volume of the swelling solvent (in cm3∙mol–1), φ2m is the volume fraction of the cross-linked polymer in the swollen gel polymer, φ2r is the polymer volume fraction in the relaxed state, i.e. after cross-linking but before swelling, f is the cross-linking functionality χ is the Flory-Huggins interaction parameter, between the polymer and solvent. The value of Mc determined from swelling data will be represented as (Mc)s. Another important parameter describing the hydrogel network properties is the effective cross-link density, ve which can be calculated as: ve = ρ p / M c (15)

The solubility parameter (δ) is an important indicator to express the interaction between a polymer and a solvent. The solubility parameter, δp, of a given poly(HEMA-co‑AAm) hydrogel was taken as that of the solvent mixture that gives the maximum swelling degree for the network. To this end mixtures of various compositions of water (δ1 = 23.40 (cal cm–3)1/2) and dimethyl formamide (δ2 = 12.13 (cal cm–3)1/2) were prepared. The solubility parameter of the mixture (δm) was calculated using the formula: xV δ +x V δ δm =1 1 1 2 2 2 x1V1 + x2V2

The equilibrium swelling of samples M1, M3, M4 and M5 in the various solvent mixtures prepared is represented in Figure 6 as a function of the solubility parameter δm. The value of δm corresponding to maximum swelling gives the solubility parameter, δp, of the respective copolymer. These are listed in Table 4.

The value of ve determined from swelling data will be represented as (ve)s. The Flory-Huggins interaction parameter χ between water and the hydrogel can be expressed in

(13)

Where xi, VI and δi are the molar fraction, molar volume, and solubility parameter of component i, respectively (V1 = 18.02 cm–3 mol–1). Table 3. Values of the enthalpy of mixing ΔHmfor polyHEMA, poly AAM and M1, M4 and M5 copolymers. Δ Hm polyHEMA (kJ∙mol–1) Temp < 35 °C –3.2

–3.5

0.74

0.00

2.1

Temp > 35 °C

–0.38

23.3

36.6

27.1

552

0.32

polyAAm

Figure 6. Experimental values of maximum swelling of the HEMA/AAm copolymers hydrogels M1, M3, M4 and M5 as a function of the solubility parameter of the solvent mixture. Polímeros , 25(6), 547-555, 2015

Synthesis and characterization of pH and temperature responsive poly(2-hydroxyethyl methacrylate-co-acrylamide) hydrogels terms of the solubility parameters of the polymer (δp) and the solvent (δ1): = χ

V1 (δ p − δ1 )2 RT

+ 0.34 (16)

where R is the gas constant and T is the absolute temperature[30]. The interaction parameter χ is also frequently calculated from the simple equation: 1 ϕ χ ap ≅ + 2m (17) 2 3

Equation 17 is derived from Equation 14 assuming that Mc goes to infinity[15]. The values of the solubility parameters and the parameter χ of the poly(HEMA-co-AAm) hydrogels are listed in Table 4. The great similarity between the values of χ obtained by Equations 16 and 17 is apparent. Table 4 illustrates that there is a clear dependence of both, the solubility parameter and χ on the copolymer composition. As the copolymer becomes richer in the more hydrophilic component, AAm, δp increases. The increase in δp causes a decrease in the polymer-solvent interaction parameter χ indicating a greater water compatibility of the copolymer hydrogels richer in AAm. This is in agreement with the composition dependence of the equilibrium water uptake values of poly(HEMA-co-AAm) hydrogels shown in Table 2. The χ values reported for the poly(HEMA‑co‑AAm) hydrogels lie in between those reported for polyAAm (0.44‑0.495)[31]and polyHEMA (0.77-0.83)[32]. Mahmudi et al.[15] studied the effect of cross-linker composition on molecular parameters of a poly(HEMA-co-AAm) hydrogel obtained by γ-irradiaton with a 6.6 kGy dose using methylenebisacrylamide as cross-linker. They reported χ values of 0.56-0.57, almost independent of the amount of cross-linker used. This value is higher than those reported in Table 4. However, the monomer composition of their reaction mixture (HEMA/AAm/H2O ≈ 25:25:50) and the radiation dose were different from those of the present work (see Table 1). In addition they do not report the temperature of experiments, which precludes comparison.

were calculated from slope of the initial linear segment of the stress-strain curves. It was then possible to calculate Mc (Equation 6) after evaluating φm using Equation 7. It has been stated that simple compression analyses can be used for the determination of effective cross-linking density without needing polymer-solvent based parameters as in the case of swelling[33]. The effective cross-linking density (νe) of the poly(HEMA-co-AAm) hyrogels were obtained from compression-strain measurements combining Equations 6 and 15: ve =

1/3 Eϕ−2m (18) 3RT

The network parameters of hydrogels obtained from mechanical measurements are listed in Table 5. They exhibit the same tendency as the values obtained from swelling measurements, in that υe decreases and Mc increases with increasing the acrylamide content of hydrogels. A very good agreement is also found between the values of the Flory-Huggins interaction parameters χ evaluated from solubility parameters (Equation 16) and those obtained from mechanical measurements. However, the values of molecular weight between cross-links, Mc, calculated from swelling experiments are about 3 to 4 times higher than those obtained from mechanical measurements. This discrepancy has been obtained before for other acrylamide based copolymers[15], and it has been explained in terms of the high sensitivity of the denominator in Equation 14 to the value of χ for highly swelled gels (φ2M<< 1). Therefore, even though the Mc values obtained from Equation 14 allow knowing the correct tendency of the variation of the molecular weight between cross-links with the copolymer composition, the more reliable Mc values should be the ones obtained from mechanical measurements.

3.5 Mechanical behavior of hydrogels The effect of composition on the mechanical properties of poly(HEMA-co-AAm) hydrogels was investigated by uniaxial compresion tests in order to evaluate the true Mc value, the interaction parameter χ and the effective cross-link density, ve. The subscript “m” will be used here to indicate that these parameters are calculated from mechanical measurements. Typical stress–strain curves of hydrogels swelled to equilibrium are given in Figure 7. As it can be seen in the figure, there is a marked dependence of the mechanical response on hydrogels composition. The magnitude of stress increased with increasing HEMA content in the copolymer hydrogel for a given strain. Elasticity modulus (E) values of hydrogels

Figure 7. Strain-stress curves of poly(HEMA-co-AAm) hydrogels.

Table 4. Solubility parameters and Flory-Huggins interaction parameter and network parameters of poly(HEMA-co-AAm) hydrogels. Copolymer

δp (cal∙cm–3)1/2

(χ)s

(χ)ap

M1 M3 M4 M5

12.93 13.63 15.06 18.96

0.54 0.52 0.51 0.50

0.537 0.522 0.511 0.508

Polímeros, 25(6), 547-555, 2015

(Mc)s ×10–5 (g∙mol–1) 3.154 2.655 2.539 1.817

(ve)s × 105 (mol∙cm–3) 1.27 1.51 2.17 3.51

553

Rapado, M., & Peniche, C. Table 5. Network parameters from stress-strain measurements. Sample

E (kPa)

pHEMA M1 M2 M3 M4 M5 pAAm

1.776 0.861 0.558 0.431 0.297 0.176 0.174

(ve)M ×105

(Mc)M ×10–5

(mol∙cm ) 1.03 1.06 1.07 1.71 2.05 2.93 4.49

(g∙mol–1) 1.250 1.063 1.032 0.664 0.523 0.434 0.253

–3

4. Conclusions Temperature and pH sensitive poly(HEMA-co-AAm) hydrogel networks were synthesized by γ-irradiation of aqueous solutions of acrylamide and 2-hydroxyethyl methacrylate at a 10 kGy dose. FTIR spectra indicate the formation of hydrogen bonding between the –OH groups of HEMA and the –NH2 groups of AAm.Water diffusion to poly(HEMA-co-AAm) hydrogels is non-Fickean. Swelling is sensitive to pH and temperature and increases with increasing the AAm content in the hydrogel. At pH values above 6 swelling of copolymers increases considerable with increasing pH due to partial hydrolysis of the acrylamide moieties. The temperature dependence of swelling exhibits a marked change around 30-35 °C that becomes more prominent for copolymers richer in AAm, which were attributed to the weakening of the hydrogen bonding between the –OH groups of HEMA and the –NH2 of AAm. The solubility parameter of copolymers, the cross‑link density of copolymers and the elastic modulus are also dependent of the AAm content of the hydrogels. It can be concluded that swelling and mechanical properties of poly(HEMA-co-AAm) copolymers prepared by γ-irradiation can be tailored to meet specific requirements for applications in different fields requiring environmentally responsive hydrogels such as medicine, pharmacy and bioengeneering.

5. Acknowledgments The authors thanks Prof. Januz M. Rosiak, for his valuable help, discussions and remarks checking the manuscript. Part of the work was done in the framework of fellowships at the Institute of Applied Radiation Chemistry, Technical University of Lotdz, Poland. We also wish to thank Lic. Rosalba ZayasMarín for her help in checking the English spelling of this article.

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(Mc)s/(Mc)M

(χ)M

(χ)S - (χ)M

---2.96 ---4.00 4.85 4.19 ----

0.59 0.55 0.54 0.52 0.51 0.50 0.50

---– 0.01 ---0.00 0.00 0.00 ----

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Synthesis and characterization of pH and temperature responsive poly(2-hydroxyethyl methacrylate-co-acrylamide) hydrogels between cross-links. Nuclear Instruments and Methods in Physics Research Section B, 265(1), 375-378. http://dx.doi. org/10.1016/j.nimb.2007.09.007. 16. Iza, M., Stoianovici, G., Viora, L., Grossiord, J. L., & Couarraze, G. (1998). Hydrogels of poly(ethylene glycol): mechanical characterization and release of a model drug. Journal of Controlled Release, 52(1-2), 41-45. http://dx.doi.org/10.1016/ S0168-3659(97)00191-0. PMid:9685934. 17. Davidson, G. W. R., 3rd, & Peppas, N. A. (1986). Solute and penetrant diffusion in swellable polymers. V. Relaxation -controlled transport in P (HEMA-cp-MMA) copolymers. Journal of Controlled Release, 3(1-4), 243-258. http://dx.doi. org/10.1016/0168-3659(86)90096-9. 18. Lurie, Y. Y. (1979). Handbook of analytical chemistry. Moscow: Chemistry Press. 19. Ham, G. E. (1966). Copolymerization. In H. F. Mark, N. Gaylord, & N. M. Bikales (Ed.), Encyclopedia of polymer science and technology (Vol. 4, pp. 165-244). New York: Interscience. 20. Murugan, R., Mohan, S., & Bigotto, A. (1998). FTIR and polarised raman spectra of acrylamide and polyacrylamide. Journal of the Korean Physical Society, 32(4), 505-512. Retrieved in 13 May 2013, from http://www.kps.or.kr/jkps/ downloadPdf.asp?articleuid=%7B43E6F355-D43F-4D4EB386-6E5B29A6F648%7D 21. Bell, C. L., & Peppas, N. A. (1995). Biomedical membranes from hydrogels and interpolymer complexes. In N. A. Peppas, & R. S. Langer (Ed.), Biopolymers II (Advances in Polymer Science, Vol. 122, pp. 125-175). Berlin: Springer-Verlag. 22. Ritger, P. L., & Peppas, N. A. (1987). A simple equation for description of solute release I. Fickian and non-Fickian release from non-swellable devices in form of slabs, sphere, cylinders or discs. Journal of Controlled Release, 5(1), 23-36. http:// dx.doi.org/10.1016/0168-3659(87)90034-4. 23. Karadag, E., Saraydin, D., & Guven, O. (2001). Radiation induced superadsorbent hydrogels. Acrylamide/Itaconic acid copolymers. Macromolecular Materials and Engineering, 286(1), 34-42. http://dx.doi.org/10.1002/1439-2054(20010101)286:1<34::AIDMAME34>3.0.CO;2-J. 24. Bocourt, M., Bada, N., Acosta, N., Bucio, E., & Peniche, C. (2014). Synthesis and characterization of novel pH-sensitive

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chitosan-poly(acrylamide-co-itaconic acid) hydrogels. Polymer International, 63(9), 1715-1723. http://dx.doi.org/10.1002/ pi.4699. 25. Hu, D. S.-G., & Lin, M. T. S. (1994). Water-polymer interactions and critical phenomena of swelling in inhomogeneous poly(acrylonitrile-acrylamide-acrylic acid) gels. Polymer, 35(20), 4416-4422. http://dx.doi.org/10.1016/0032-3861(94)90101-5. 26. Wang, J., Wu, W., & Lin, Z. (2008). Kinetics and thermodynamics of the water sorption of 2-Hydroxyethyl methacrylate/styrene copolymer hydrogels. Journal of Applied Polymer Science, 109(5), 3018-3023. http://dx.doi.org/10.1002/app.28403. 27. Peniche, C., Cohen, M. E., Vázquez, B., & San Román, J. (1997). Water sorption of flexible networks based on 2-hydroxyethyl methacrylate-triethylenglycol dimethacrylate copolymers. Polymer, 38(24), 5977-5982. http://dx.doi.org/10.1016/S00323861(96)01058-0. 28. Mark, J. E., & Erman, B. (Ed.). (1988). Rubberlike elasticity: a molecular primer. New York: Wiley. 29. Flory, P. J., & Rehner, J., Jr. (1943). Statistical mechanics of swelling of crosslinked polymer networks II. The Journal of Chemical Physics, 11, 521-526. 30. Teraoka, I. (2002). Polymer solutions: an introduction to physical properties. New York: Wiley. 31. Orwol, R. A., & Chong, Y. S. (1999). Polyacrylamide. In J. E. Mark (Ed.), Polymer data handbook (pp. 247-251). Oxford: Oxford University Press. 32. Day, J. C., & Robb, I. D. (1981). Thermodynamic parameters of polyacrylamides in water. Polymer, 22(11), 1530-1533. http://dx.doi.org/10.1016/0032-3861(81)90324-4. 33. Uzun, O., Hassnisaber, M., Şen, M., & Guven, O. (2003). Enhancement and control of cross-linking of dimethylaminoethyl methacrylate irradiated at low dose rate in the presence of ethylene glycol dimethacrylate. Nuclear Instruments and Methods in Physics Research Section B, 208, 242-246. http:// dx.doi.org/10.1016/S0168-583X(03)01112-1. Received: Feb. 03, 2014 Revised: May 25, 2015 Accepted: Aug. 15, 2015

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http://dx.doi.org/10.1590/0104-1428.2158

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

Molecular weight and tacticity effect on morphological and mechanical properties of Ziegler–Natta catalyzed isotactic polypropylenes Ismael Amer1,2*, Albert van Reenen1 and Touhami Mokrani2 Department of Chemistry and Polymer Science, University of Stellenbosch, Stellenbosch, South Africa 2 Department of Civil and Chemical Engineering, College of Science, Engineering and Technology, University of South Africa – UNISA, Johannesburg, South Africa

*ismaelamer77@yahoo.co.za

Abstract The morphological and mechanical properties of Ziegler–Natta catalyzed isotactic polypropylenes as influenced by the molecular weight and tacticity was investigated. Polypropylene samples were injection moulded into standard disks with a HAAKE MiniJet II injection moulder at 190 °C and 200 bar for morphological and mechanical tests. The morphological and mechanical properties of specimens were investigated by means of optical microscope (OM), scanning electron microscopy (SEM), microhardness (MH) and dynamic mechanical analysis (DMA). The samples exhibited a typical α-modification spherulite structure of isotactic polypropylenes crystallized from the melt. It was found that the most important factor affecting the structure and properties of these polymers is the isotacticity content. A clear molecular weight effect was also found for samples with low molecular weights. The microhardness and storage modulus values increased as crystallinity did. Accordingly, isotacticity degree is considered as the main parameter affecting the crystallinity of samples. Keywords: mechanical property, molecular weight, morphological property, polypropylene, tacticity.

1. Introduction Polypropylene is one of the most common polymers in use today. Its good mechanical properties and relatively low price result in the continuous growth of its production and the expansion of its market. Its continuously increasing application accelerates research in all related fields, including the preparation of isotactic polypropylene based composites and blends[1]. The mechanical and physical properties of polypropylene are influenced by a number of factors. The mechanical properties of the majority of polypropylene homopolymers are, apart from processing conditions, influenced by their rheological and crystallization behaviour. The degree of crystallinity is considered to be the most influential property affecting the physical and mechanical properties of a polypropylene sample[2-4]. An increase in crystallinity is often related to an increase in properties such as the stiffness or storage modulus of a sample, while other factors such as the impact strength generally decrease with increasing crystallinity. The storage modulus can be defined as the ratio of stress to strain under vibratory conditions which can be calculated from data obtained from either free or forced vibration tests, in shear, compression, or elongation. The Stiffness can be defined as the rigidity of the material to which it resists deformation in response to an applied force. An increase in crystallinity can lead to an increase in the lamellar thickness which leads to higher storage modulus and stiffiness values[2-4]. The effects of molecular weight[5-8], molecular weight distribution[9] and tacticity[10-13] on the crystallization have been investigated by several authors. Cheng et al.[6] showed that the linear growth rate of crystals decreases with the increase of molecular weight,

556

but the overall crystallization rate might increase because an increasing number of intramolecular folded-chain nuclei could result in a higher nucleation density[7]. For samples with similar molecular weights and different tacticities the linear crystal growth rate might increase by three orders of magnitude when the isotacticity (mmmm %) of isotactic polypropylene increases from 78.7 to 98.8%[11,14]. The degree of crystallinity of isotactic polypropylene is commonly in the range of 40 to 70%[1]. Atactic polypropylene, on the other hand, is considered uncrystallizable, since the chain structure lacks regularity. Isotactic polypropylene can crystallize in three different crystal forms as was described by Bruckner et al.[15], depending on the polymer structure and the crystallization conditions: the α-form with a monoclinic, the γ-form with an orthorhombic and the β-form with a hexagonal unit cell[15]. The molecular weight has also been shown to influence the glass transition temperature (Tg) of polymers, with higher molecular weight samples having a higher Tg[16]. This in turn influences the mobility of chains at room temperature, and since polypropylene has a Tg range in the region of 0 °C, variations in the Tg temperature range can have an effect on the ability of the material to displace energy at low temperatures. The aim of the current study is to explore narrowly the structure–property relationships of polypropylene. The morphological and mechanical properties of different samples were investigated by means of optical microscope (OM), scanning electron microscopy (SEM), microhardness

Polímeros , 25(6), 556-563, 2015

Molecular weight and tacticity effect on morphological and mechanical properties of Ziegler–Natta catalyzed isotactic polypropylenes (MH) and dynamic mechanical analysis (DMA) to demonstrate the effect of isotacticity and molecular weight on the properties of polypropylene samples. To the best of our knowledge the investigation of how the molecular weight affect the thermal, morphological and mechanical properties of isotactic polypropylene has not been done yet to polymer studies which has previously been overlooked.

2. Experimental Section 2.1 Materials Polypropylene polymers and fractions used in this study were selected from our previous study as shown in Amer and van Reenen[17] and Table 1. These polymers were prepared using a commercial Ziegler-Natta catalyst with 2.78 wt% Ti content purchased from Star Chemicals & Catalysts Co. (China) and labeled as P3−P17. P4(120), P5(120), P8(120) and P9(110) refer to polypropylene TREF fractions eluted at 120 and 110 °C, respectively.

2.2 Polymerization procedure All polymerization reactions were carried out under an inert gas atmosphere. The polymerization reactions were carried out in a 350-mL stainless steel Parr autoclave with a gas inlet and pressure gauge. Typically the reactor was charged with the catalyst (43 mg, Ti content 2.78 wt%) and triethylaluminium (2 mmol, Al/Ti mole ratio 80) in toluene (25 mL). The catalyst solution was stirred for 5 min and then the propylene was added. The reactor was pressurized with hydrogen and the contents stirred for 1 h at room temperature. The reaction was then quenched by the addition of 100 ml 10% HCl/MeOH. The resulting polymer was filtered off, washed several times with methanol, and subsequently dried under vacuum at 80 °C for 15 h, to yield about 3-5 g of polypropylene as a white powder.

2.3 Sample preparation for mechanical tests Test specimens were injection moulded into standard disks for morphological and mechanical tests with a HAAKE MiniJet II injection moulder. The injection moulding temperature was 190 °C and the injection pressure was 200 bar. The dimensions of the standard disks are 20.0 mm in diameter and 1.5 mm in thickness.

2.4 Preparation of etching reagent Permanganic etching of polyolefins was used to prepare samples for the study of the morphology. This technique has been used in several studies involving polyolefins[18-21]. Potassium permanganate (1 g) (obtained from Sigma‑Aldrich) was dissolved in 100 mL of a concentrated solution of 33 vol % phosphoric acid and 67 vol % sulphuric acid (Sigma‑Aldrich). The solution was prepared by adding potassium permanganate very slowly to the beaker containing both acids, with rapid agitation. After adding all the potassium permanganate, the beaker was closed and the content stirred until all the potassium permanganate was dissolved (a dark green purple solution formed). All polypropylene samples were etched at room temperature.

2.5 Etching procedure Specimens from each polymer, with approximate dimensions of length 10 mm, width 5 mm and thickness 1.5 mm, were cut from the disks prepared by injection moulding. Each sample was immersed in about 10 mL of the etching reagent in a beaker for a period of 60 minutes. This permanganic acid solution preferentially etches the amorphous part of the polymer in the spherulites in such a way that the lamellae then appear clearly. Subsequently, the specimens were carefully washed with hydrogen peroxide, distilled water and acetone, in order to avoid any artefacts caused by pollution effects. Samples were finally dried in a vacuum oven at 45 °C for 5 hr.

2.6 Polymer characterization A Zeiss Axiolab OM, (magnification × 50-100 μm) with a high resolution camera CCD-IRIS (Sony) was used to examine the etched piece, to investigate the crystal structure. SEM analysis of etched piece was performed using a Leo® 1430VP scanning electron microscope operated at 15 kV of acceleration voltage at room temperature. All the surfaces to be studied were coated with gold under vacuum in order to eliminate any undesirable charge effects during the SEM observations. Samples for compressive DMA were analyzed using a Perkin Elmer DMA 7e calibrated according to standard procedures. The samples were first melted at 180 °C for 8 minutes and then melt pressed at 5 MPa and same temperature. The samples were analyzed using a 5 °C/min

Table 1. Characterization data of the polypropylenes and fractions. Samples

Mwa (g/mol)

Mw/Mn

mmmmb (%)

Tmc (°C)

Tcc (°C)

ΔHmc (J/g)

Xc (%)

P3 P4 P5 P8 P14 P17 P4(120) P5(120) P8(120) P9 (110)

184 759 252 956 312 580 228 960 215 397 65 498 195 693 207 823 178 423 110 387

6.1 5.4 4.1 6.4 5.9 8.2 4.3 2.9 3.5 3.4

93.0 94.0 96.0 94.0 86.0 93.0 98.0 96.0 96.0 98.0

161.2 160.6 161.9 162.0 157.5 156.5 160.0 161.0 158.6 158.7

124.2 116.5 118.4 124.4 119.8 120.8 116.0 118.0 116.3 118.4

104.5 103.9 108.9 103.9 90.5 100.5 119.7 103.2 106.9 110.7

50.0 50.0 52.0 50.0 43.0 48.0 57.0 49.0 51.2 53.0

determined by GPC. bdetermined by NMR. cmeasured by DSC.

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Amer, I., van Reenen, A., & Mokrani, T. heating ramp with an applied force oscillating at a frequency of 1 Hz. The static force was kept constant at 110% of the dynamic force. The temperature range analyzed was between –40 °C and 230 °C. MH measurements were conducted on a UHL microhardness tester equipped with a Vickers indenter. Measurements were obtained using an indentation speed of 25 μm/s and a dwell time of 15 s. Samples were analyzed at indentation loads of 10 gf. Ten measurements were recorded for each sample analyzed.

3. Results and Discussions 3.1 Morphological properties Table 1 summarizes the polypropylene polymers and fractions used in this study. 3.1.1 Optical Microscopy analysis First, the effect that the molecular weight has on the crystal structure of different polypropylene samples was studied. Figure 1a-c shows OM micrographs of polypropylene fractions P5(120), P4(120) and P9(110) respectively, which differ in molecular weight (see Table 1). They exhibit a typical α-modification spherulite structure of isotactic polypropylenes crystallized from the melt. These micrographs show that all the isotactic polypropylene fractions have well-defined and large α-spherulitic morphology. The spherulites grew,

impinged on each other, and formed particular polygonal spherulites with clear boundaries. Indications are that, since all the observed spherulites grew at the same rate and their observed size can be considered uniform, the nuclei are formed immediately after cooling to the crystallization temperature and their number remains constant thereafter. The only effect of the molecular weight that can be noticed in Figure 1 is the slight morphological differences in the sign of birefringence, magnitude of the birefringence and spherulite texture. Second, the effect of tacticity on the crystal structure of different isotactic polypropylenes was studied. Figure 2a-c show OM micrographs of polypropylenes P5, P4 and P14, which differ in their tacticities (Table 1). The graphs show that, under similar crystallization conditions, the dimensions of the crystal structures of these different isotactic polypropylene samples decrease in size with tacticity. This effect can be explained by the restriction of movement of polymer chains caused by chain defects in low tacticity polymers during the crystallization process, resulting in slower crystallization and hence the formation of smaller spherulites. Similar results were obtained in other studies[22-24]. 3.1.2 Scanning Electron Microscopy analysis Figure 3a-c shows SEM micrographs of the typical crystallization morphologies of the isotactic polypropylene fractions P5(120), P4(120) and P9(110) respectively, which differ in molecular weight. All these isotactic polypropylenes

Figure 1. Optical micrographs of isotactic polypropylene fractions: (a) P5(120) (Mw = 207823 g/mol), (b) P4(120) (Mw = 195693 g/mol) and (c) P9(110) (Mw =110387 g/mol) (500x magnification).

Figure 2. Optical micrographs of isotactic polypropylene polymers: (a) P5 (mmmm = 96%), (b) P4 (mmmm = 94%) and (c) P14 (mmmm = 86%) (500x magnification). 558

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Molecular weight and tacticity effect on morphological and mechanical properties of Ziegler–Natta catalyzed isotactic polypropylenes revealed well-defined and large spherulitic morphology, comprising a mixture of α1 (disordered) and α2 (ordered) crystal form structures. The spherulites grew, impinged on each other, and formed particular polygonal spherulites with clear boundaries. Moreover, one can clearly see the individual lamellae and lamellar branching structure in the SEM micrographs in Figure 3. The samples consist of crosshatch-type lamellar branching structures, which is the typical characteristic of the α crystal form of isotactic polypropylenes[25,26]. In contrast to OM results, clear differences can be distinguished between the three different isotactic polypropylene samples shown in Figure 3. These differences exist in the variety of spherulite sizes and spherulite types classified by their appearance, including the sign and nature of birefringence and crystal lattice. The average diameter of P5(120) spherulites is about 5-15 μm (Figure 3a). Smaller dominant α spherulites (about 5-10 μm) are observed for P4(120) and P9(110) (Figure 3b, c respectively). Figure 4a-c illustrates SEM micrographs of polypropylenes (P5, P4 and P14) that differ in their tacticities (Table 1). Similar to those results obtained from OM above, SEM also shows in Figure 4, that the sizes of the spherulites were decreased drastically with decreasing tacticity. In addition, with decreasing tacticity, the spherulites showed less perfection and the sharp spherulite boundaries became more diffuse (Figure 4c). The sample P5 with 96.0% tacticity has the biggest spherulite sizes (15-25 μm) while samples P4 and P14 with 94.0% and 86.0% tacticities have spherulite sizes about 5-15 and 1-3 μm respectively. SEM micrographs (Figure 4a, b) also show small dimples on the etched surface

of samples P5 and P4 may grow to craters and holes. This is due to the extractions of the rubbery materials by the etchant solution[19,20,27,28].

3.2 Mechanical properties In order to correlate the structure of the polypropylene polymers with the mechanical properties, the samples were analyzed using microhardness and DMA. 3.2.1 Microhardness According to literature[29-33], all the parameters that lead to an increase of crystallinity and crystallite sizes (lamellar thickness) will also lead to higher MH values. Hence, the higher the isotacticity, the greater the MH values obtained. The effect of the molecular weight and molecular weight distribution on the crystallinity of the polypropylene samples P3, P4, P5, P8, P17, P4(120) and P8(120) was investigated and results are illustrated in Figure 5. The samples of low molecular weight generally have a broader molecular weight distribution, and vice versa. The molecular weight distribution has an effect on the crystallinity: the samples with a higher degree of crystallinity have a lower molecular weight distribution, and samples with lower degree of crystallinity have a higher molecular weight distribution. On the other hand, there is a slight increase in the crystallinity of the samples with an increase in the molecular weight. Figure 6 shows the combined effect of the molecular weight and isotacticity on the degree of crystallinity of the polypropylene samples P3, P4, P5, P8, P17, P4(120), P8(120) and P9(110). There is a remarkable increase in crystallinity as the isotacticity is increased, from about 93% to 98%.

Figure 3. SEM micrographs of isotactic polypropylene fractions: (a) P5(120) (Mw = 207823 g/mol), (b) P4(120) (Mw = 195693 g/mol) and (c) P9(110) (Mw =110387 g/mol) (3000x magnification).

Figure 4. SEM micrographs of isotactic polypropylene polymers: (a) P5 (mmmm = 96%), (b) P4 (mmmm = 94%) and (c) P14 (mmmm = 86%) (3000× magnification). Polímeros, 25(6), 556-563, 2015

559

Amer, I., van Reenen, A., & Mokrani, T. In general, the samples with high molecular weight have high isotacticity, as expected, since the more stereospecific sites have a higher propagation constant rate (Kp). This is in agreement with the results obtained by Sakurai et al.[34] with regards to the relationship between isotacticity and molecular weight. Moreover, Figures 5 and 6 show that the crystallinity of the samples was largely affected by the isotacticity, which clearly dominates over other effects such as molecular weight and molecular weight distributions. Nevertheless, when the isotacticity values from the polymer samples are similar, the molecular weight and molecular weight distributions exert significant influence on crystallinity.

the crystallinity of these polypropylene polymers is the isotacticity content, the combined effect of the molecular weight and isotacticity on the MH of the samples P3, P4, P5, P8, P17, P4(120), P8(120) and P9(110) is illustrated in Figure 8. It appears that there is a significant increase in the MH with increasing isotacticities of the samples. Hence, it can be said that the higher the isotacticity the greater the MH. This means that the most important factors affecting the MH of these polypropylene polymers are those that lead to an increase in crystallinity. In addition, the main parameter which affects the crystallinity can be considered to be the degree of isotacticity of the samples.

Looking at the combined effect of the molecular weight and crystallinity on the microhardness of the samples P3, P4, P5, P8, P17, P4(120), P8(120) and P9(110), as shown in Figure 7, generally one can see that there is a major increase in the MH with increasing molecular weight and crystallinity of the samples. Since that the most important factor affecting

It is reasonable that the higher isotacticity content allows easier recrystallization upon the application of an external force to the sample, thus improving the hardness of the sample upon indentation. The magnitude of the effect of the tacticity of the polypropylenes on the properties of the polymer has also been discussed by De Rosa et al.[35].

Figure 5. The effect of molecular weight and molecular weight distribution on the crystallinity of the isotactic polypropylene polymers P3, P4, P5, P8, P17, P4(120) and P8(120).

Figure 7. The combined effect of molecular weight and crystallinity on the microhardness of the isotactic polypropylene polymers P3, P4, P5, P8, P17, P4(120), P8(120) and P9(110).

Figure 6. The combined effect of molecular weight and isotacticity on the crystallinity of the isotactic polypropylene polymers P3, P4, P5, P8, P17, P4(120), P8(120) and P9(110).

Figure 8. The combined effect of molecular weight and isotacticity on the microhardness of the isotactic polypropylene polymers P3, P4, P5, P8, P17, P4(120), P8(120) and P9(110).

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Polímeros , 25(6), 556-563, 2015

Molecular weight and tacticity effect on morphological and mechanical properties of Ziegler–Natta catalyzed isotactic polypropylenes 3.2.2 Dynamic Mechanical analysis 3.2.2.1 Effect of molecular weight Figure 9 shows the storage modulus of different isotactic polypropylene samples with different molecular weights, as a function in temperature. The storage modulus values increase with increasing molecular weight in the temperature range measured for the different samples. This is in agreement with the results obtained above from the microhardness test. Similar behaviour has been observed by other researchers[8]. The reason for the increase in the storage modulus values with increasing molecular weight is due to the higher degree of crystallinity and the presence of a larger number of molecular weight entanglements per chain for the higher molecular weight polymers. Furthermore, an increase in the lamellar thickness as the molecular weight increases also leads to higher storage modulus values[8]. The detailed plot of loss tangent (tan δ) of these isotactic polypropylene samples, as a function of the temperature, ranging from –40 to 80 °C is presented in Figure 10. The tan δ curves represent the ratio of the ability of the material to store and lose energy, which is sometimes referred to as the clamping ability of a material. It can also be taken as a measurement of the impact properties of the material. It is apparent from Figure 10 that β-transition, corresponding to the Tg of isotactic polypropylenes, which occurs over the temperature range 10-20 °C, is slightly shifted to a higher

Figure 9. Storage modulus curves as a function of temperature for isotactic polypropylene samples with various molecular weights.

Figure 10. Tan δ curves as a function of temperature for isotactic polypropylene samples with various molecular weights. Polímeros, 25(6), 556-563, 2015

temperature as the molecular weight increases (from 16 °C for P3 sample with Mw 184 759 g/mol to 20 °C for P5 sample with Mw 312 580 g/mol). The samples with lower molecular weights are less crystalline and therefore contain more amorphous material. The explanation is that the chains have far greater mobility in the amorphous phase in the lower molecular weight samples compared to the samples with higher molecular weights. We do, however, also have to take into consideration the change in the molecular packing in the amorphous phase. Denser packing of the molecular chains leads to a reduction in the molecular motion. The areas of the β-transitions of the samples, after subtraction of a linear baseline, are given in Figure 11. Figure 11 shows that the magnitude of the β-transition increases with increasing molecular weight. Similar results were obtained by Stern et al.[8], who found that the higher molecular weight polymers are generally characterized by larger β-transition. In fact, a decrease in the mechanical transition of the β process is associated with a reduction in the mobility of the polymer chains in the amorphous phase[8,36,37]. 3.2.2.2 Effect of isotacticity Figure 12 shows the difference in the storage modulus curves, as a function of temperature, for isotactic polypropylene samples with various isotacticities. Compared to the results obtained from varying the molecular weight, isotactity of isotactic polypropylene samples has further effect on the overall viscoelastic response as shown in Figure 12. The storage modulus is greater in the higher isotactic polypropylene samples than in those with lower tacticity, over the whole temperature range studied (1.51 × 108 Pa for P4(120) with 98 mmmm % vs 0.84 × 108 Pa for P14 with 86 mmmm %) as shown in Figure 12. Moreover, this effect is also observed in the location and intensity of the β-transition temperature, as shown in Figure 13. As the isotactic content increases, the location of the β-transition temperature is considerably shifted to higher temperatures, but its intensity decreases significantly in the higher isotactic polypropylene fractions P5(120) and P4(120), as can be seen in Figure 14, which shows the areas of the β-transitions of the different samples. All of these features

Figure 11. The magnitude of the area of the β-transition for isotactic polypropylene samples with various molecular weights. 561

Amer, I., van Reenen, A., & Mokrani, T. can be associated with the higher degree of crystallinity that more regular chains can be reached, i.e. as isotacticity is increased in the isotactic polypropylene macromolecules. Therefore, the lowest content of amorphous regions is in the P4(120) sample, because its higher crystallinity (Xc = 57%, Table 1) leads to its higher storage modulus, a decrease in magnitude of the β-transition and the displacement of its location to higher temperatures due to the higher hindrance of motions within the crystalline phase.

Figure 12. Storage modulus curves as a function of temperature for isotactic polypropylene samples with various isotacticities.

4. Conclusions The relationship between structure and properties was established for different polypropylene samples through their morphological and mechanical characterizations. The effect of isotacticity and molecular weight on the properties of polypropylene samples was investigated. The most important factors affecting the structure and properties of these polypropylene samples are those that lead to an increase of crystallinity. Consequently, the main parameter is the degree of isotacticity, followed by molecular weight. OM and SEM results showed that all isotactic polypropylene samples had well-defined α-spherulitic morphology. OM and SEM also showed that tacticity had a greater effect on the morphological structure of the isotactic polypropylenes than molecular weight. A decrease in isotacticity leads to a clear decrease in the dimensions of the crystal structures for the different isotactic polypropylene samples. Results of the MH and DMA showed that all the parameters that lead to an increase in crystallinity and crystallite sizes (lamellar thickness) will provide higher MH, storage modulus and β-transition temperature values. The crystallinity of the samples was shown to be affected by the molecular weight and molecular weight distribution, as well as by the isotacticity of the samples.

5. Acknowledgements The authors would like to thank 1) the International Centre of Macromolecules and Materials Science (Libya) for financial support for Ismael Amer and 2) Mrs Sonja Brandt for the English editing.

6. References

Figure 13. Tan δ curves as a function of temperature of isotactic polypropylene samples with various isotacticities.

Figure 14. The magnitude of the area of the β-transition of isotactic polypropylene samples with various isotacticities. 562

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http://dx.doi.org/10.1590/0104-1428.2047

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

Design of conformal cooling for plastic injection moulding by heat transfer simulation Sabrina Marques1*, Adriano Fagali de Souza2, Jackson Miranda3 and Ihar Yadroitsau4 Additive Manufacturing Laboratory, SENAI Institute of Innovation in Manufacturing Systems, Joinville, SC, Brazil 2 Advanced Manufacturing Systems Laboratory, Universidade Federal de Santa Catarina – UFSC, Joinville, SC, Brazil 3 Tupy University Center, Sociedade Educacional de Santa Catarina – SOCIESC, Joinville, SC, Brazil 4 Department of Mechanical and Mechatronic Engineering, Central University of Technology, Bloemfontein, South Africa 1

*sabrina.marques@sc.senai.br

Abstract The cooling channels of a mold for plastic injection have to be as close as possible to the part geometry in order to ensure fast and homogeneous cooling. However, conventional methods to manufacture cooling channels (drilling) can only produce linear holes. Selective laser melting (SLM) is an additive manufacturing technique capable to manufacture complex cooling channels (known as conformal cooling). Nevertheless, because of the high costs of SLM the benefits of conformal collings are still not clear. The current work investigates two designs of conformal coolings: i) parallel circuit; ii) serial circuit. Both coolings are evaluated against to traditional cooling circuits (linear channels) by CAE simulation to produce parts of polypropylene. The results show that if the conformal cooling is not properly designed it cannot provide reasonable results. The deformation of the product can be reduced significantly after injection but the cycle time reduced not more than 6%. Keywords: plastic injection moulding process, conformal cooling, additive manufacturing.

1. Introduction The industries of plastic parts are a driving force for the current market. Such manufacturing process requires tooling known as moulds to manufacturing plastic parts and products. The mould is the most important component in the process of manufacturing a plastic part because it influences the cycle time and the quality of the product. Optimal properties of plastics parts can be achieved only when the correct mould temperature is used and maintained during the manufacturing process. The mould temperature influences the: mechanical properties; shrinkage; warpage; surface quality; cycle time and the flow length for thin walled parts[1,2]. The cooling time during an injection moulding process usually represents about 2/3 of the total cycle time[3]. Therefore, any reduction on the cooling time will have a great repercussion on the complete production time. The efficiency of the cooling circuit directly influences the quality and the cooling time of the part. At least 60% of visible defects such as wrapping recorded in the injected component may be related to the inefficiency of the cooling system[4]. The cooling system inside the mould cavity is manufactured by drilling machining. It means only linear channels are possible to manufacture. Then, in most of the cases a uniform heat transfer is not obtained. Researches can be found in literature about linear cooling circuits[5-7]. New technologies to manufacture metal parts have been developed during the past decades known as Addictive

564

Manufacturing techniques (AM). Among them the selective laser melting (SLM) is unique powder-based technologies that produce objects from metal powders with complex geometries. The mechanical properties are comparable to that one of tool steel[8]. Using SLM techniques the cooling channels can be manufactured following the product topography. These complex channels are kwon as conformal cooling. An expressive reduction of the production time of a plastic part together to a better product property is expected with the conformal cooling channels[9,10]. However, the cost of this technology is still very high. Thus, the benefits of the conformal cooling for a plastic injection mould should be determined before the mould manufacture using SLM. According to Hamdy et al.[11] to achieves an optimum thermal reduction and shrinkage rate distribution throughout the product, the conformal cooling system layout must be optimized. According to Dalgarno and Stewart[12] the use of conformal cooling channels in an injection mould can result in a significant reduction in the cycle time. Ilyas et al.[13] stated an improvement of the productivity and energy saving due to the high efficiency of the heat transfer by the conformal cooling channels. Hsu et al.[14] identify the efficiency of conformal cooling by three-dimensional simulation to reducing the cooling time displacement. However, none of these have investigated the design of conformal cooling and options of the cooling channels.

Polímeros , 25(6), 564-574, 2015

Design of conformal cooling for plastic injection moulding by heat transfer simulation Focusing this issue, the current works aims to understand the behavior of different conformal cooling designs. Using the geometry of a plastic part as a workpiece, two designs of conformal cooling are proposed: parallel circuit (manifold) and serial circuit. Both conformal cooling circuits are evaluated against the traditional cooling circuits (linear channels) by CAE simulation to produce parts of polypropylene.

Table 1. Type of coolant flow[1,2].

1.1 Thermal analysis

1.3 Temperature variation of polymer melt

After the injection phase, the heat transferred to the mould cavity by the molten material needs to be extracted, ensuring the solidification of the polymer melt. After the polymer is solidified, it is ejected from the mould cavities. Generally, the recommended moulding conditions for injecting polypropylene (based on the standard ISO 1873‑2:2007[15]) suggests the mould temperature about 40°C and the melt temperature was 200°C[16]. Heat exchange occurs between the mould and the molten material, causing a constant temperature increase of the mould up to a temperature where it is stabilized. According to Park and Dang[17], when the heat balance is established the heat flux supplied to the mould and the heat flux removed from the mould are in equilibrium. The heat balance is expressed by Equation 1.

All flow problems involve solving the equations of conservation of mass, momentum and energy, these equations are described as Kennedy[19] and Li and Shen[20]:

Qm + Qc + Qe = 0 (1)

where Qm, Qc and Qe (W/m2) are respectively: the heat flux the melt, the heat flux exchange with coolant and environment.

1.2 Cooling time The cooling time is proportional to the square of the thickest wall of the part and the largest runner diameter powered of 1.6, and inversely proportional to the thermal diffusivity of the polymer melt. These relationships are given by the Equation 2[1,2].

(

)

Tc ∝ Thw2 + Dr1.6 ∝

1 (2) a

Reynolds number, (Re) 4 000 < (Re) 2 300 < (Re) < 4 000 100 < (Re) < 2 300 (Re) < 100

Mass equation: ∂ρ + ( ∇.ρv ) = 0 (5) ∂t

Momentum ∂v ρ = − ∇.ρ + ∇.ηγ  − ρ [ v.∇v ] equation: ∂t

Energy equation:

ρ.C p (

ρCv

Re =

ρUd (4) η

where ρ is the density of the coolant (kg/m3), is the averaged velocity of the coolant (m/s), is the diameter of the cooling channel (m), and η is the dynamic viscosity of the coolant (kg m-1s-2 ). The type of coolant flow can be determined by the Reynolds number Re, as listed in Table 1. Polímeros, 25(6), 564-574, 2015

(6)

∂T ∂p + v.∇T ) =βT ( + v.∇p) + ∂t ∂t (7)

ηγ 2 + ∇(k ∇T )

where ρ is the material density (kg/m3), t is the time (s), is the melt velocity (m/s), η is the viscosity (Pa.s), is the shear rate (s-1), Cp is the specific heat (J/K), β is the coefficient of thermal expansion (K-1), k is the thermal conductivity (W/m.K), and T is the temperature (K). Adopting a system of Cartesian plane and assuming the thin thickness of the cavity compared with other dimensions, the mass and momentum equations can be reduced to Equations 8 and 9[19,20]. ∂  ∂p  ∂  ∂p  S + S  = ∂x  ∂x  ∂y  ∂y  1 H  ∂p β − ∫ K 2 − H  ∂t ρC p

 2 ∂  ∂T ηγ ∂z  K ∂z  

2   z      ∫ hh +− dz    2 h + η   1 z  where Tc (s) is the cooling time, Thw (m) is the thickness at = S dz −   ∫  2 h − η h+ 1 the thickest part of the part wall, Dr (m) is the diameter of  ∫ h − dz   η    the largest runner, and is the thermal diffusivity (m2/s). Thus,    doubling the wall thickness quadruples the cooling time. 2   z      ∫ hh +− dz    The thermal diffusivity of polymer melt is defined 2 + h η   1 z  S dz −   = according to the Equation 3[18]. ∫  2 h − η h+ 1  ∫ h − dz   K η    (3) a=   

where ρ is the thermal conductivity (W/m.K), ρρ is the density (kg/m3), and Cv is the specific heat constant volume (J/kg.K). The Reynolds number is described according to the Equation 4 [1,2].

Type of flow Turbulent Transition Laminar Stagnated

(8)

      dz   

(9)

where h+ and h- are respectively the highest and lowest z coordinate of the frozen layer’s position, K is the coefficient of expansion (K-1), and represents half the wall thickness (m). Assuming that convection in the z direction can be ignored, the energy equation is described according to Equation 10[19,20].  ∂T  ∂t +   ∂T ρC p  vx ∂x  ∂ v T  y  ∂t

   ∂T ∂  ∂T   +  = βT + ηγ 2 +  k  (10) ∂y ∂z  ∂z      565

Marques, S., Souza, A. F., Miranda, J., & Yadroitsau, I. The left side of the energy equation represents the rate of temperature change and convection, while the term on the right describes the expansion/compression by heating, viscous dissipation and heat conduction to the mould However, the Hele-Shaw model, neglects the velocity and pressure variation in the thickness direction, resulting in a two-dimensional flow with a heat flow problem in the plane of flow and an additional problem of thermal conduction in the thickness direction[4]. Consequently, this technique is not recommended for models with high wall thickness and complex geometries, such as used in this study. For these cases a simulation analysis using full 3D technology is required. By matching Equations 11 and 12 he variation of the viscosity (η, Pa.s) as a function of the shearing rate (γ, s-1), temperature (T, K) and pressure (p, Pa) is obtained.

190mm (L) × 155mm (W) × 65mm (H). The dimensions of bottom insert are 190mm (L) × 155mm (W) × 61mm (H). Braskem H201 polypropylene was used as the plastic material for injection moulding simulation cording to Table 3. The Mouldflow V10 CAE software was used for the simulation and the specific parameters were assumed to be: ▪ Material of the inserts: AISI P20. ▪ Injection temperature: 230°C. ▪ Injection time: 1 sec. ▪ Coolant fluid: water. ▪ Maximum pressure injection machine: 180 MPa. ▪ Maximum force closure machine: 7000 ton.

η0 η ( γ, T , p ) = (11)  η0.γ  1+  *   τ 

Three designs of the cooling channels are proposed: i) series conformal cooling; ii) parallel conformal cooling; iii) linear cooling channels (ordinary) as presented ahead.

 − A . (T − D2 − D3.P )  η0 (T , p ) = D1.exp  1  (12) A2 + T − D2  

2.1 Conformal cooling design

The coefficients τ*, ɳ , D1, D2, D3, A1 and A2 are adjusted dada. τ* is the tension (Pa) on the transition of the shearing behavior, D2 temperature, (K) glass transition (Tg).

2. Materials and Methods This work investigates the designs of the conformal cooling circuits of a plastic injection mould. A plastic tray of eggs-holder for refrigerators was used as the work piece. Table 2 shows the work piece, inserts and the injection mould used in this study. The work piece geometry has five equidistant cavities interconnected by a thickness of 2 mm and 140 mm overall diameter. The supply channel dimensions are: 60 mm in length, 6.5 mm diameter at the entrance and draft angle of 2°. The dimensions of top insert are

Indentifying the requirements for conformal cooling designing was the first task in this phase. The position of the channels, the pitch distance, its diameter and its length must be considered[3,22]. Based on this information, these dimensions were obtained from[23], as presented by Table 4 and Figure 1. Considering the moulded product (work piece) is 2mm thickness, the dimensions of the cooling channels is assumed to be: b = 5mm. Thus: a = 2.5 × 5 = 12.5mm and c = 1.7 × 5 = 8.5mm. But besides dimensions of the circuits, the topology of the conformal cooling must be designed. Because this is a new approach to design cooling circuits, a guide about its geometry was not found in literature. Therefore this work investigates the conformal cooling topography. It is proposed and evaluated by FEA-CAE simulation the design of conformal cooling in series channels and conformal cooling in parallel channels, as detailed ahead.

Table 2. Work piece, inserts and mould design[21]. Product

Inserts Top

Sketch of the Mould

Bottom

566

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Design of conformal cooling for plastic injection moulding by heat transfer simulation 2.1.1 Conformal cooling in series channels In this case the conformal cooling channels were designed so that the cooling liquid passes through each cavity one at a time, in series. Series cooling channels are connected in a single loop from the coolant inlet to its outlet (Table 5). This type of cooling channel network is the most commonly used in practice (for linear channels). By design, if the cooling channels are uniform in size, the

coolant can maintain its turbulent flow rate through its entire length. Turbulent flow enables the heat to be transferred more effectively. According to Park and Dang[17], for linear channels it is not appropriate use channels in series circuits in the following situations: ▪ If the length of the series circuit requires a higher pressure than the pump capacity can support. ▪ The physical constraints in the mould design means that the mould cannot be effectively cooled with a series circuit.

2.1.2 Conformal cooling in parallel channels

Figure 1. Optimal design of a ‘three dimensional channel system[23]. Table 3. Physical and mechanical properties typical of polypropylene. Properties Fluidity index (g/10min) Density (g/cm3) Secant flexural modulus 1% (GPa) Tensile strength in the flow (MPa) Stretching in the flow (%) Rockwell Hardness (R scale) Izod impact strength 23°C (J/m) Heat deflection temperature 0,455 MPa (°C) Heat deflection temperature 1,820 MPa (°C) Vicat softening temperature 10 N (°C)

PP H 201 20 0.905 1.5 34 12 102 23 97 57 154

Parallel cooling channels are straight drilled channels in which the coolant flows from a supply manifold to a collection manifold. The collection manifold is designed with a larger diameter than the cavity’s channels (Ø12 mm). Due to the flow characteristics of the parallel cooling channels, the flow rate along various cooling channels may be different, depending on the flow resistance of each individual cooling channel. The varying of the flow rate, cause the heat transfer efficiency of the cooling channels to vary from one to another. As a result, cooling of the mould may not be uniform with a parallel cooling channel configuration, but a balanced parallel circuit provides uniform heat extraction. However, it is suggested only to use a parallel circuit if the model has one of the following circumstances[17]: ▪ The pressure drop over a series circuit is too high to be realistic. ▪ An area of the mould cannot be effectively cooled with a series circuit.

Table 6 shows the proposed design of the parallel conformal cooling circuit.

Table 4. Optimal design of a three dimensional channel system[23]. Wall thickness of moulded product (mm) 0-2 >2-4 >4-6

Hole diameter, b (mm) 4-8 >8-12 >12-14

Centreline distance between holes, a (mm) 2-3×b 2-3×b 2-3×b

Distance between centre of holes and cavity, c (mm) 1.5-2×b 1.5-2×b 1.5-2×b

Table 5. Design of the series conformal cooling channels. Top insert

Bottom insert

Sketch view

Series conformal cooling channels

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Marques, S., Souza, A. F., Miranda, J., & Yadroitsau, I. 2.2 Design of the ordinary cooling system with linear channels

3. Results and Discussions

The linear cooling system manufactured by drilling usually requires plugs to connect the channels and to close the circuit. Because of the constraints of this process, the circuits developed normally do not follow the geometry of the product, as they are linear channels. TableÂ 7 shows the design of the linear cooling channels proposed in this work.

The simulation results are presented by graphics comparing the three cooling systems investigated. Some graphics had to be presented in an individual scales, in order to have a good representation of the cooling phenomenon. FigureÂ 2 shows the circuit pressure inside the cooling circuit generated from a cooling analysis to show the

Table 6. Design of the parallel conformal cooling channels. Descriptions

Top insert

Bottom insert

Sketch view

Parallel conformal cooling channels

Table 7. Design of the series cooling channels. Descriptions

Top insert

Bottom insert

Linear cooling channels

Figure 2. Circuit pressure results 568

PolĂmeros , 25(6), 564-574, 2015

Design of conformal cooling for plastic injection moulding by heat transfer simulation distribution of pressure along the cooling circuits. The circuit pressure is responsible for making the refrigerant fluid circular in the channels. The value of the circuit pressure is directly dependent on the geometry of the circuit. It is one of the important factors to evaluate whether a cooling system is viable, since high pressures are not preferred[24]. The linear circuit operates with a lower pressure flow, because it is a simple geometry with little restriction to flow. It is observed that both conformal cooling circuits

required higher pressures to operate and have larger pressure drops along the circuit because of the high level of flow restriction (complex geometry). Figure 3 and 4 shows the flow rate of the coolant and Reynolds number inside the cooling circuit. The circuit’s flow rate and the circuit’s Reynolds number are used to determine the flow rate required to achieve a turbulent coolant flow. The flow rate itself is not the dominant factor in heat extraction, but it should

Figure 3. Circuit flow rate result.

Figure 4. Reynolds number result. Polímeros, 25(6), 564-574, 2015

569

Marques, S., Souza, A. F., Miranda, J., & Yadroitsau, I. be the minimum requirement to achieve the necessary Reynolds number. The flow rate is constant for a series circuit, but not for a parallel circuit. A Reynolds number of 4,000 or higher represents a more turbulent flow, which is preferred for cooling applications. However, the higher the Reynolds number in the circuit, the more energy is required to pump it through the circuit. Hence, the ideal Reynolds number for cooling circuits is 10,000. The pumping losses associated with a Reynolds number higher than 10,000 outweigh the heat transfer gains that can be achieved with higher Reynolds numbers[1,2]. The simulations shown in Figure 3a, b and Figure 4a, b shows that for linear and for series circuits the flow rate and the Reynolds number are constant. Whereas for the parallel circuit the flow rate and the Reynolds number varies (Figure 3c and Figure 4c). The flow rate for linear circuit and for series circuit is sufficient to achieve turbulent coolant flow and their values are close to 10,000, which is the optimal value for the Reynolds number. Whereas in most regions along the parallel circuit the flow rates are insufficient to achieve turbulent coolant flow. And in the regions where the flow rates are sufficient to achieve turbulent coolant flow, the values are very high compared to the optimal value of Reynolds number. Figure 5 shows the coolant temperature results inside the cooling circuit. The temperature of the coolant is related to the flow rate and pressure applied in the circuit. The coolant temperature varies along the circuit. It can be noticed that on the serial circuits, this variation occurred linearly and the temperature at the output channel is higher than the temperature along the circuit, unlike a

parallel circuit, where the highest temperature occurs at the centre of the circuit. The difference between the in and outlet coolant temperatures should be no more than 3°C. Higher values may indicate a high mould surface temperature, and should be avoided in order to get a good product quality[1,2]. The difference of temperature observed on Figure 5 is according to the specified threshold, but the conformal cooling circuits presented higher temperature variations due to the higher heat transfer, because the channels are closer to the mould cavity, compared with conventional circuits. The lowest variation of the temperature along the circuit between the conformal cooling circuits was achieved by the series conformal cooling circuit (Figure 5b). Figure 6 identifies that the regions with the highest temperature is the interior of the cavities of the top insert and the centre of the bottom insert. This helps the further analysis of the heat removal of the mould. Focusing the investigation on the regions with a greater concentration of heat (Figure 6), Figure 7 shows the efficiency of the heat extraction according to the regions of the mould for each cooling system designed. Value close to 1 indicates a more efficient region of the circuits. Figure 7b shows that the series circuit had the highest efficiency (up to 1) in the warm regions of the mould cavities. Compared to some regions, where the efficiency is zero in regions with a high heat load. The most appropriate cooling design the heat removal efficiency must be high in the regions of the product and also be homogeneous to keep the residual stress as uniform as possible. Figure 8 indicate the mould surface temperature inside the mould, during the cycle.

Figure 5. Circuit coolant temperature result. 570

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Design of conformal cooling for plastic injection moulding by heat transfer simulation Figure 8b shows the series cooling circuit had the greatest efficiency and propitiated a homogeneous cooling around the mould. The most regions of the part kept temperatures around 28°C (blue colour in the graph). Satisfactory performance of the cooling system requires a homogeneous temperature in this interface without any hot spots, which is the great cause of tensions and warpages in parts. When the interface is cooled homogeneously with small temperature differences, the chance of defect is less and the quality of the part will be better[1,2]. Figure 9 shows the time required to reach the part ejection temperature. The series conformal cooling channels have the shortest cycle time (Figure 9b). The difference in cycle time (ejection temperature) between the series

circuit and the conventional circuit was 1.33 sec., which means a 6% reduction. The simulation indicates there is not a significant reduction in the cycle time when using conformal cooling. Probably it happens because either the product geometry is not so complex with deeper regions more difficult the cool down, or because the injection channel is extracted together with the product. Therefore, the geometry of the product and a hot runner injection must be taken into account to design a conformal cooling and identify the worth case. The deflection of the product after the injection was also simulated according to the cooling design as illustrated by Figure 10. According to simulation, the linear cooling system caused a deflection about 0.29 mm inside the product cavities and in its centre. A relative large area of the product was significantly deformed. The best results can be found on the series cooling system (Figure 10b). Only small area inside the product’s cavities suffered small deformations. The mould with parallel cooling channels propitiates deformations inside the product’s cavities, similar to the serial channels case, but further deformation can be observed on the product border, about 0.1mm. As previously stated by Figure 8b the mould with a series cooling circuits results in a homogeneous cooling around the mould. Now, Figure 10b shows that this homogeneous cooling of the mould propitiates the lowest deflection compared with its counterparts.

Figure 6. Regions of the mould with a greater concentration of heat.

Table 8 presents the summary of the simulations results of the three different cooling systems.

Figure 7. Circuit heat removal efficiency result. Polímeros, 25(6), 564-574, 2015

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Marques, S., Souza, A. F., Miranda, J., & Yadroitsau, I.

Figure 8. Mould temperature result.

Figure 9. Time to reach ejection temperature result.

Figure 10. Deflection of the injected part. 572

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Design of conformal cooling for plastic injection moulding by heat transfer simulation Table 8. Summary results. Cooling channels Circuit pressure (psi) Top insert Bottom insert Circuit flow rate (l/min) Top insert Bottom insert Circuit Reynolds number Top insert Bottom insert Circuit coolant temperature (°C) Top insert Bottom insert Circuit heat removal efficiency Top insert Bottom insert Mould temperature (°C) Time to reach ejection (s) Part deflection

Ordinary

Series

Parallel

(a)

(b)

(c)

1.36-0.16

14.5-0.0014

14.24-1.46

1.36-0.16

14.5-0.0072

14.06-0.62

6.11

1.49

29.89-0.01

6.11

2.82

38.5-0.03

18029

7049

88254-49

18029

13308

75776-146

25-25

25-26

25-28

25-25

25-27

0.96-0.45

0.98-0.0

0.5-0.0

0.64-0.31

0.98-0.0

0.5-0.0

25.0-36.2 23.07 Higher

28.0-30.0 21.74 Smaller

4. Conclusions Conformal cooling designed for plastic injection moulding could be an attractive alternative to improve the plastic product quality, reducing cycle time and energy consumption. However due to the high costs to manufacture such cooling channels (addictive manufactures techniques must be applied) together to the lack of knowledge about its quantitative benefits, this application is still insipient. Concerning this issue the current work proposed two designs for conformal cooling. By CAE simulation (Moldflow V10) both designs were evaluated against ordinary cooling channels (linear ones) in a case study. The performance of these 3 cooling designs were accessed by the simulation and discussed. The remarkable points are: 1. The proposed conformal cooling design in series channels proved to be the best option for the case investigated. The reduction of the product deformations was significant if compared with the liner channels and parallel conformal cooling. However the reduction of the cycle time was not as expected. It propitiates 6% reduction to produce a part. Probably it happens because either the geometry used as work piece didn’t have a high degree of complexity or because the injection channel is extracted with the product. 2. Conformal cooling tends to be worthy when the product has deep regions which linear cooling channels cannot archive reasonable hot transfer. Identify this point is the challenge for the mould designers. Current commercial CAD/CAE software does not offer a wizard routine to add this identification. 3. The flow rates for the series circuit are sufficient to achieve turbulent coolant flow and their values are close to 10,000, which is the optimal value for the Polímeros, 25(6), 564-574, 2015

25.0-33.0 23.19 Intermediate

Reynolds number. It propitiates a very homogeneous mould temperature and then an injected product with the smaller deformation. 4. The parallel conformal cooling propitiates an insufficient flow rate to achieve turbulent coolant flow in most area of the circuit. In the regions where the turbulent coolant flow was possible to be archived, the values are very high compared to the optimal value of Reynolds number.

5. Acknowledgements The authors thank CAPES, CNPq and FAPESC for supporting this research project and also thank the partners SOCIESC, DIPI-ENISE and Villares Metals.

6. Reference 1. Duleba, B., & Greskovic, F. (2012). Conformal cooling for plastics injection moulding. Strojar, 1-5. Retrieved in 02 August 2014, from http://www.it-strojar.sk/articles/00016.pdf 2. Dimla, D. E., Camilotto, M., & Miani, F. (2005). Design and optimization of conformal cooling channels in injection moulding tools. Journal of Materials Processing Technology, 164-165, 1294-1300. http://dx.doi.org/10.1016/j.jmatprotec.2005.02.162. 3. Ching, Z. L., & Chou, M. H. (2002). Design of the cooling channels in nonrectangular plastic flat injection mold. Journal of Manufacturing Systems, 21(3), 107-186. http://dx.doi. org/10.1016/S0278-6125(02)80160-1. 4. Zhou, H., Yan, B., & Zhang, Y. (2008). 3D filling simulation of injection moulding based on the PG method. Journal of Materials Processing Technology, 204(1-3), 475-480. http:// dx.doi.org/10.1016/j.jmatprotec.2008.03.017. 5. Li, C. L. (2001). A feature-based approach to injection mould cooling system design. Computer Aided Design, 33(14), 10731090. http://dx.doi.org/10.1016/S0010-4485(00)00144-5. 573

Marques, S., Souza, A. F., Miranda, J., & Yadroitsau, I. 6. Li, C. L., Li, C. G., & Mok, A. C. K. (2005). Automatic layout design of plastic injection mould cooling system. Computer Aided Design, 37(7), 645-662. http://dx.doi.org/10.1016/j. cad.2004.08.003. 7. Li, C. G., & Li, C. L. (2008). Plastic injection mould cooling system design by the configuration space method. Computer Aided Design, 40(3), 334-349. http://dx.doi.org/10.1016/j. cad.2007.11.010. 8. Yadroitsev, I., & Smurov, I. (2011). Surface morphology in selective laser melting of metal powders. Physics Procedia, 12(part A), 264-270. http://dx.doi.org/10.1016/j.phpro.2011.03.034 9. Wang, Y., Yu, K.-M., Wang, C. C. L., & Zhang, Y. (2011). Automatic design of conformal cooling circuits for rapid tooling. Computer Aided Design, 43(8), 1001-1010. http:// dx.doi.org/10.1016/j.cad.2011.04.011. 10. Dimitrov, D., & Moammer, A. (2010). Investigation towards the impact of conformal cooling on the performance of injection moulds for the packaging industry. Journal for New Generation Sciences, 8(1), 29-46. Retrieved in 14 June 2014, from http:// reference.sabinet.co.za/webx/access/electronic_journals/ newgen/newgen_v8_n1_a3.pdf 11. Hassan, H., Regnier, N., Pujos, C., Arquis, E., & Defaye, G. (2010). Modeling the effect of cooling system on the shrinkage and temperature of the polymer by injection moulding. Applied Thermal Engineering, 30(13), 1547-1557. http://dx.doi. org/10.1016/j.applthermaleng.2010.02.025. 12. Dalgarno, K. W., & Stewart, T. D. (2001). Manufacture of production injection mould tooling incorporating conformal cooling channels via indirect selective laser sintering. Proceedings of the Institution of Mechanical Engineers. Part B, Journal of Engineering Manufacture, 215(10), 1323-1332. http://dx.doi. org/10.1243/0954405011519042. 13. Ilyas, I., Taylor, C., Dalgarno, K., & Gosden, J. (2010). Design and manufacture of injection mould tool inserts produced using indirect SLS and machining processes. Rapid Prototyping Journal, 16(6), 429-440. http://dx.doi.org/10.1108/13552541011083353. 14. Hsu, F. H., Wang, H. C., Huang, C. T., & Chang, R. Y. (2013). Investigation on conformal cooling system design in injection moulding. Advances in Production Engineering & Management, 8(2), 107-115. http://dx.doi.org/10.14743/apem2013.2.158. 15. International Organization for Standardization – ISO. (2007). ISO 1873-2: plastics – polypropylene (pp) moulding and extrusion materials – part 2: preparation of test specimens and determination of properties. Geneva: ISO.

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16. Daiyan, H., Andreassen, E., Grytten, F., Lyngstad, O. V., Luksepp, T., & Osnes, H. (2010). Low-velocity impact response of injection-moulded polypropylene plates – Part 1: Effects of plate thickness, impact velocity and temperature. Polymer Testing, 29(6), 648-657. http://dx.doi.org/10.1016/j. polymertesting.2010.05.003. 17. Park, H. S., & Dang, X. P. (2012). Design and simulation-based optimization of cooling channels for plastic injection mould. In C. Volosencu (Ed.), New technologies: trends, innovations and research (pp. 19-44). Rijeka: InTech. Retrieved in 16 June 2014, from http://cdn.intechopen.com/pdfs-wm/34669.pdf 18. Venerus, D. C., Schieber, J. D., Iddir, H., Guzman, J. D., & Broerman, A. W. (1999). Relaxation of anisotropic thermal diffusivity in a polymer melt following step shear strain. Physical Review Letters, 82(2), 366-369. http://dx.doi.org/10.1103/ PhysRevLett.82.366. 19. Kennedy, P. (1999). CAD, CAM, & CAE. Lexington: Mouldflow Corporation. 20. Li, C. S., & Shen, Y. K. (1995). Optimum design of runner system balancing in injection moulding. International Communications in Heat and Mass Transfer, 22(2), 179-188. http://dx.doi.org/10.1016/0735-1933(95)00003-8. 21. Marques, S., Souza, A. F., Miranda, J. R., & Santos, R. F. F. (2014). Evaluating the conformal cooling system in moulds for plastic injection by CAE simulation. In Proceedings of the 9th International Conference on Industrial Tools and Material Processing Technologies. Ljubljana: Slovenian Tool and Die Development Centre. Retrieved in 14 June 2014, from http:// www.met.uni-miskolc.hu/refbase/_publication_files/2014/273_ FeldeI.2014_Estimationofheattransfer.pdf 22. Xu, X., Sachs, E., & Allen, S. (2001). The design of conformal cooling channels in injection molding tooling. Polymer Engineering and Science, 41(7), 1265-1279. http://dx.doi. org/10.1002/pen.10827. 23. Siegfried, M. (2009). Optimized mould temperature control procedure using DMLS. EOS Whitepaper, p. 1-10. Retrieved in 02 April 2013, from http://www.compositesworld.com/ uploadedFiles/Publications/MMS/Articles/Internal/EOS_ WP_DMLS2_ENG_12.pdf. 24. Rees, H. (1995). Mould engineering. Munich: Hanser. Received: Jan. 21, 2015 Revised: July 07, 2015 Accepted: Aug. 06, 2015

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http://dx.doi.org/10.1590/0104-1428.2057

Synthesis and photostabilizing performance of a polymeric HALS based on 1,2,2,6,6-pentamethylpiperidine and vinyl acetate Marcelo Aparecido Chinelatto1*, José Augusto Marcondes Agnelli2 and Sebastião Vicente Canevarolo2 Departamento de Engenharia de Materiais, Universidade de São Paulo – USP, São Carlos, SP, Brasil Departamento de Engenharia de Materiais, Universidade Federal de São Carlos – UFSCar, São Carlos, SP, Brasil 1

*mchinelatto@sc.usp.br

Abstract Polymeric hindered amine light stabilizers (polymeric HALS) have been extensively studied because they combine a high ability to protect the polymers against harmful effects of weathering with minimum physical loss. In this study a new polymeric N-methylated HALS was synthesized by the radical copolymerization of a cyclic tertiary amine with vinyl acetate (VAc). 4-Acryloyloxy-1,2,2,6,6-pentamethylpiperidine (APP), the cyclic tertiary amine, was prepared by the initial conversion of 2,2,6,6-tetramethyl-4-piperidinol derivatives via two different routes. The APP/VAc copolymer synthesized was characterized by size exclusion chromatography (SEC), Fourier transform infrared spectroscopy (FTIR) and carbon-13 nuclear magnetic resonance (13C NMR). The photostabilizing performance, particularly the induction period of polypropylene (PP) films containing different concentrations of APP/VAc copolymer, when exposed to accelerated aging, was comparable to that of PP films compounded with commercial polymeric HALS. Keywords: degradation, HALS, hindered amines, photo-oxidation, photostabilizer.

1. Introduction Low stability against heat and ultraviolet (UV) radiation are considered the main limitations of polyolefins and hamper their use in several applications. Particularly when exposed to UV radiation and in the presence of oxygen, polyolefins undergo a photo-oxidative process, whose intensity can range from a superficial change, damaging the aesthetics of the product, to significant structural modifications, affecting their physical and mechanical properties[1,2]. The addition of hindered amine light stabilizers (HALS) to a polymer is the most effective way to stabilize it against UV radiation. HALS derivatives of 2,2,6,6-tetramethylpiperidine are blended into a large number of commercial polymers via extrusion process[3,4]. However the high temperatures and high shear rates used during the polymer extrusion may cause the volatilization of the low molecular weight HALS, significantly reducing their concentration and so their protection against photo-oxidation[5]. Besides volatilization, other physical phenomena, such as compatibility with polymeric matrix and resistance to extraction by water and organic solvents during the end use also contribute to the reduction of the low molecular weight photostabilizers efficiency[6-8]. The synthesis of HALS of higher molecular weight by addition or condensation copolymerizations[9-11] or the grafting of photostabilizers onto the preformed polymer[12-16] are usual strategies described by the literature to eliminate or minimize such undesirable effects. The preparation of polymeric HALS by the copolymerization of hindered amines with ethylene or propylene is also an alternative route to reduce their physical loss[17,18]. Auer et al.[19] described a novel and facile one-step synthetic route to prepare a HALS and its copolymerization

Polímeros, 25(6), 575-580, 2015

with ethylene using a metallocene/methylalumoxane (MAO) as a catalyst system. In a previous paper we studied the functionalization of 2,2,6,6-tetramethyl-4-piperidinol, a hindered amine by methyl acrylate and its subsequent radical copolymerization with vinyl acetate (VAc) and styrene (Sty)[20]. During the copolymerization the hydrogen atom bonded to the nitrogen in the piperidine ring was labile and reacted with propagating radicals. The high hindrance imposed by the four methyl groups on the radical formed in the nitrogen atom inhibited the formation of a new chemical bond with monomers whose pendant group has a large volume, as the case of Sty. The difficulty of establishing an effective chemical bond between nitrogen and Sty resulted in the formation of a soluble copolymer. However the pendant group in VAc would not be large enough to avoid the formation of an effective chemical bond whose radical is generated in the nitrogen atom, forming the crosslinked structure. The undesired reactivity of the hydrogen bonded to the nitrogen in the hindered amine during the radical copolymerization could be overcome by the alkylation of N-H HALS. The photostabilizing effect of tertiary HALS, like N-alkyl HALS in polypropylene (PP) or polyethylene (PE) is comparable or superior to that of the parent secondary HALS[21-23]. The mechanism by which N-methyl HALS inhibits a polymer oxidation involves the attack of N-alkyl groups of tertiary HALS by oxidative active species, leading to its dealkylation and producing parent secondary HALS and carboxylic acid. The oxidation of this secondary amine leads to a stable nitroxyl radical. However other possible photostabilization mechanisms are the quenching

575

T T T T T T T T T T T T T T T T T T T T

Chinelatto, M. A., Agnelli, J. A. M., & Canevarolo, S. V. of excited oxygen-polymer charge transfer complexes and hydroperoxides deactivation (CTC)[21,24,25].

composition of the copolymer were determined according to the procedure previously described[20].

In this paper the conversion of 2,2,6,6-tetramethyl-4piperidinol into APP was studied so as to enable its subsequent radical copolymerization with a vinyl monomer. The APP was prepared by two different routes involving the alkylation and functionalization of 2,2,6,6-tetramethyl-4-piperinol. The APP and intermediate compounds generated during the conversion of the piperidinol structure were characterized by FTIR and 13C NMR. Subsequently a new polymeric HALS was synthesized via the radical copolymerization of APP with VAc and characterized by FTIR, 13C NMR and SEC and its photostabilizing effectiveness was evaluated in PP films.

The molecular weight distribution curve of the APP/VAc copolymer was determined by SEC on a Waters 510 Chromatograph with a refractive index detector using three Ultrastyragel linear columns. The calibration curves were plotted using twelve samples of monodisperse polystyrene standards, ranging from 480 to 1x106 g mol-1. Tetrahydrofuran was used as the solvent at a flow rate of 1.0 mL min-1 at room temperature. Molecular weights were determined by Milennium 2010 software.

2. Materials and Methods 2.1 Materials 2,2,6,6-Tetramethyl-4-piperidinol (Aldrich, 98%), acryloyl chloride (Aldrich, 96%), triethylamine (Aldrich, 99%), methyl acrylate (Aldrich, 99%), methyl iodide (Aldrich 99%) and titanium (IV) isopropoxide - Tipox (Aldrich, 97%) were used as received. Vinyl acetate (Aldrich, 98%) was used after distillation. Azo-bis-isobutyronitrile (AIBN), used as a thermal decomposition initiator, was kindly donated by Bayer from Brazil. An isotactic PP was provided by Braskem S.A. (Brazil). According to the manufacturer, at a melt flow rate of 15 g 10 min-1 (230oC/2.16 Kg) this PP is considered a medium melt flow rate homopolymer with a general-purpose additive package, suitable for injection molding and fiber extrusion. The efficiency of the new polymeric HALS synthesized here to protect the PP films against photo-oxidation was compared with that of a commercial product Chimassorb 944 (HALS-1) – poly[[6-[(1,1,3,3-tetramethylbutyl)amino]1,3,5-triazine-2,4-diyl][(2,2,6,6-tetramethyl-4-piperidinyl) imino]-1,6-hexanediyl[(2,2,6,6-tetramethyl-4-piperidinyl) imino]] from Basf (Germany).

2.2 Characterization The APP and the intermediate compounds generated during the alkylation and functionalization of piperidinol were characterized by FTIR using a Spectrum 1000 Spectrophotometer from Perkin-Elmer, at a 4000-400 cm-1 wavenumber range, with 24 scans for each spectra and 2 cm-1 resolution. The samples were dissolved in chloroform, cast over potassium bromide (KBr) windows and dried under an infrared lamp. The 13C NMR spectra of 2,2,6,6-tetramethyl-4-piperidinol derivatives and APP/VAc copolymer were recorded on a Varian Unityplus 400 spectrometer operating at a 100.57 MHz for the carbon-13 nucleus using a 5 mm probe at room temperature. For quantitative measurements, the sample solutions (250 mg/250 μL) were prepared in CDCl3. The experiments were performed in inverse gated decoupling mode. The parameter for these studies were 90o pulse width 10 μs, relaxation delay 15 s, acquisition time 0,68 s and 3.000 repetitive scans were taken for a good signal-to-noise ratio. The molar ratio of the monomers and the chemical 576

2.3 Preparation of APP The preparation of APP was performed by two different routes. In the first the APP was prepared by the alkylation of 4-acryloyloxy-2,2,6,6-tretamethylpiperidine (4ATP). Thus, firstly 4-acryloyloxy-2,2,6,6-tretamethylpiperidine (4ATP) was prepared by a transesterification reaction, according to the procedure previously described[20,26]. 4ATP (0.5 g; 2.4 mmol), methyl iodide (2.8 mL; 45.2 mmol) and potassium carbonate (0.28 g; 2.0 mmol) were added in a single-neck flask reactor with a mixture of distilled water (1.87 mL) and chloroform (4.69 mL) according to the procedure described by Kurumada et al.[23]. The mixture was maintained under stirring for one week at room temperature. It was then poured into an aqueous ammonia solution and extracted with chloroform. The chloroform was dried with anhydrous potassium carbonate and evaporated under reduced pressure. The global yield of the APP prepared by this route was 5.5%. In the second route the alkylation of 2,2,6,6-tetramethyl4-piperidinol (0.5 g; 3.2 mmol) was carried out by the same method described above and resulted in the formation of 1,2,2,6,6-pentamethyl-4-piperidinol (PPOl), a crystalline solid that cannot be polymerized by free radicals. In the functionalization of this tertiary cyclic amine, PPOl reacted with acryloyl chloride and triethylamine according to a procedure described in the literature[27]. The liquid obtained at the end of the reaction was also the APP. The final yield of APP was 12.5%, being this higher yield the main advantage when the two routes are compared. IR data APP (cm-1): C=O 1733; C=C 1636; O-C(=O)-C 1295; C-O 1064; N-CH3 958.

13 C NMR APP (ppm): N-CH3 27.7; C-CH3 32.8; >C=O 165.6.

2.4 Synthesis of APP/VAc copolymer The APP/VAc copolymer, a polymeric N-methylated HALS, was prepared by free radical copolymerization, using toluene as the solvent. The AIBN concentration was 1.0 wt% by weight of the monomer mixture (3.61x10-2 mmol), the molar ratio of APP to VAc was 1:6, and the mass ratio of the solvent to the monomers mixture was 2.5:1. The radical copolymerization was performed in a 250 mL flask containing the above-mentioned proportions of comonomers and toluene, under nitrogen atmosphere. The radical copolymerization of APP with VAc was carried out for 10 hours at 60oC and the polymer was reprecipitated in petroleum ether, collected and dried for 8 hours under vacuum at 40oC. Polímeros , 25(6), 575-580, 2015

Synthesis and photostabilizing performance of a polymeric HALS based on 1,2,2,6,6-pentamethylpiperidine and vinyl acetate 2.5 Performance of APP/VAc copolymer in accelerated weathering 2.5.1 Preparation of the test samples The APP/VAc copolymer and a HALS-1 were dispersed in the PP by a hot dissolution process. A 250 mL glass beaker with 5.0 g PP and 50 mL of xylene was placed on a hot plate at 100oC and maintained at this temperature under stirring until the complete dissolution of the PP. Thereafter the respective amounts of HALS were added to the polymer solution. The heating and manual stirring were maintained until the complete evaporation of the solvent and the light-stabilized polymers were dried in a vacuum oven for 24 hours at 70oC. The samples were prepared with different concentrations of polymeric HALS in such a way that the equivalent piperidine functionality concentration kept constant. A comparison based on the equivalent functional concentration, i.e., equal nitrogen content of piperidine ring enables the proper evaluation of the HALS effectiveness with different molecular weights and structures[28]. Therefore the samples were prepared with the equivalent piperidine functionality concentrations of 0.093 g/Kg PP and 0.140 g /Kg PP. PP films for FTIR spectral characterization were prepared by compression molding using a laboratory press set at 190oC, 10 MPa pressure and compression time of 2 min. Films for the accelerated weathering were selected having transmittance of 35±2% at 1170 cm-1. The absorption at 1170 cm-1 is ascribed to C-C stretch and proportional to the film thickness. 2.5.2 Accelerated weathering exposure The photo-oxidation was performed using an Atlas Weather-Ometer model Ci65 equipped exclusively with xenon lamps, daylight filter and irradiance at 340 nm of 0.35W/m2.nm. The exposure conditions were 102 min light at 63oC (black panel temperature) and 18 min light and water spray. A new sample was withdrawn every 100 hours of exposure and did not return to the Weather-O-Meter. The photo-oxidation of the PP films was monitored by an FTIR spectrophotometer (Spectrum 1000, Perkin-Elmer) under the same conditions described above. To eliminate the influence of the film thickness on the absorption of the carbonyl groups, the normalized carbonyl index (CI) was calculated as the ratio of the absorption intensity at 1713 cm-1, due to the C=O stretch, and 1170 cm-1 due to the C-C stretch (CI=A1713/A1170). The evolution of the degradation process was evaluated by measuring the variation of the carbonyl index (CIt-CI0), where CIt is the normalized carbonyl index determined at time t and CI0 is the normalized carbonyl index before aging (t=0).

polydispersity is 1.56. The molecular weight in the range from 2,000 to 3,000 g mol-1 is suitable to guarantee that HALS will not leach out of the polymer[29]. Stabilizers whose molecular weight is higher than this range are usually more resistant to volatility or leaching out, but the increase in the molecular weight decreases their solubility and compatibility with the polymeric matrix. Gugumus studied the effect of molecular weight of polyacrylate HALS on the stabilization efficiency of the cast PP films and observed that the major drop in performance occurs if the molecular weight increases from 2,700 to 6,800 g mol-1[30]. The results also showed that even for very large molecular weights there remains some contribution of the HALS to UV stability, which increases with the increase in the HALS’ concentration. The APP/VAc copolymer was characterized by 13C NMR and its spectrum is shown in Figure 1. The signals at 170.3 ppm and 174.2 ppm are due to the carbonyl carbons. The signal at 170.3 ppm is assigned to the carbonyl carbon of the VAc because it is more shielded than the carbonyl carbon of the APP. The opening of the double C=C bond of the APP during the copolymer synthesis causes a 165.6 ppm to 174.2 ppm displacement of the signal relative to the carbonyl carbon. The signals at 27.7 ppm and 32.9 ppm are related to carbons of the methyl groups bonded to carbon and nitrogen in the APP unit, respectively. The chemical shift at 18.7 ppm is assigned to the carbon of the methyl group of the VAc. The molar ratio of APP to VAc in the copolymer was determined by the ratio between the area of each sign and the number of nuclei responsible for the signal. Integrating the signal related to the carbons of methyl groups present in the chemical structures of the APP and VAc and dividing them by the number of nuclei responsible for each signal yielded a 1:2.5 molar ratio of APP to VAc in the copolymer. This molar ratio is higher than that of APP:VAc in the feed, i.e. 1:6, which indicates that the APP is more reactive than the VAc under polymerization conditions. Knowing the weight-average molecular weight ( M w ) value and the molar ratio of APP:VAc in the copolymer and assuming that the molecular weights of APP and VAc are 225.3 g mol-1 and 86.1 g mol-1, respectively, the APP/VAc copolymer

3. Results and Discussions 3.1 Synthesis of APP/VAc copolymer The new polymeric HALS was synthesized by radical copolymerization, employing a 1:6 molar ratio of APP to VAc in the feed. The SEC results showed that the molecular weight distribution curve is monomodal, the number-average molecular weight ( M n ) is 4,300g mol-1 and Polímeros, 25(6), 575-580, 2015

Figure 1. 13C-NMR spectrum of APP/VAc copolymer. 577

Chinelatto, M. A., Agnelli, J. A. M., & Canevarolo, S. V. is comprised of 15 units of APP and 38 units of VAc, on average. The results indicate that the polymeric HALS synthesized under this conditions is a statistical copolymer.

3.2 Performance of APP/VAc copolymer in accelerated weathering The curves of the normalized carbonyl index (CI) as a function of the aging time of the PP films with and without polymeric HALS are shown in Figure 2. According to the literature these curves can be characterized by two distinct regions: in the first, known as induction period, the value of the normalized CI is zero, whereas in the other a rapid increase in the normalized CI is monitored, following a

quasi-linear evolution[31]. Although it is clear that the curves in the Figure 2 showed a substantial dispersion of data, their behaviors were assumed to follow a linear correlation passing through the last point in the induction period. Usually studies quantifying the evolution of the degradation process in polymers, applying non-destructive characterization methods, utilize the same set of samples that are removed from the aging environment, tested and returned back to further age and test again. In experiments carried out using this procedure, particularly FTIR measuring the normalized CI do show a quasi-linear relationship. Here a different film is tested each time and so greater data dispersion is expected. One reason could be attributed to the possible fractional

Figure 2. Normalized CI as a function of the aging time of PP films – (●) control; (□) [APP/VAc] = 0.093 g/Kg PP; (∆) [HALS-1] = 0.093 g/Kg PP; (■) [APP/VAc] = 0.140 g/Kg PP and (▲) [HALS-1] = 0.140 g/Kg PP. 578

Polímeros , 25(6), 575-580, 2015

Synthesis and photostabilizing performance of a polymeric HALS based on 1,2,2,6,6-pentamethylpiperidine and vinyl acetate Table 1. Induction period and normalized CI of PP films under accelerated aging conditions. PP film

[EPF]a (x10–2g/Kg PP)

Induction Period (h) 300

CIt-CI0

Control

_______

0.6b

APP/VAc

9.3 14.0

500 900

0.8c 0.8c

HALS-1

9.3 14.0

600 600

0.6c 0.6c

Equivalent piperidine functionality concentration. bVariation in the normalized CI (CIt-CI0) after 600 h under accelerated aging. cVariation in the normalized CI (CIt-CI0) after 2000 h under accelerated aging. a

precipitation of the PP and polymeric HALS during the preparation of the films via cast from solution resulting in a less homogeneous distribution of the photostabilizer in the polymer matrix. According to Figure 2a, the PP films without photostabilizers (control) showed low stability against the photo-oxidation given an induction period of only 300 hours. After that the rate of CI increase is quite high reaching already 0.6 at 600 hours of exposure to the accelerated weathering. Longer aging times produce PP films that are too brittle and the samples crack and crash easily, precluding the normalized CI to be measured. The PP films containing APP/VAc copolymer and HALS-1 as polymeric HALS showed a significant increase in the induction period in comparison with the control films. The induction periods of the PP films containing an equivalent piperidine functionality concentration of 0.093 g/Kg-PP of APP/VAc copolymer (Figure 2b) and HALS-1 (Figure 2c) were 500 hours and 600 hours, respectively, nearly twice those exhibited by the control films. By increasing the equivalent piperidine functionality concentration of APP/VAc copolymer in PP films to 0.140 g /Kg-PP (Figure 2d), the induction period was further increased to 900 hours. However the increases in the nitrogen content of piperidine ring in the light-stabilized films with HALS-1 (Figure 2e) did not result in a longer induction period expected. The values o f induction period and normalized carbonyl index after 2000 hours of exposure to the accelerated weathering of the PP films with and without photostabilizers are shown in Table 1. The performance of the novel HALS proposed here reach values for the induction period close to the values shown by the referenced commercial compound. The protection of the new stabilizer at longer exposition times did show a lower performance, indicating its application mainly to items with shorter life cycles.

4. Conclusions A new polymeric N-methyl HALS was synthesized by the radical copolymerization of 2,2,6,6-tetramethyl4-piperidinol derivatives with VAc. The N-substituted piperidinol derivatives were prepared via alkylation reaction with methyl iodide by two different routes. The alkylation of 2,2,6,6-tetramethyl-4-piperidinol and reaction of acryloyl chloride produced a final yield of 12.5%, against 5.5% obtained by the alkylation of 4ATP, which is an important Polímeros, 25(6), 575-580, 2015

advantage of the second route. While dispersed in a PP matrix APP/VAc copolymer did protected it against the weathering environment showing similar performance in terms of induction period as a commercial polymeric HALS. Its application would include mainly protection to goods of low cycle life.

5. Acknowledgements We express our acknowledgments for the invaluable collaboration of Elói J. Esmanhoto (in memoriam), the National Council for Scientific and Technological Development (CNPq) for the financial support and the Institute of Technology for Development (LACTEC) for the accelerated weathering tests.

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Polímeros , 25(6), 575-580, 2015

http://dx.doi.org/10.1590/0104-1428.1986

Biobased additive plasticizing Polylactic acid (PLA) Mounira Maiza1*, Mohamed Tahar Benaniba1, Guilhem Quintard2 and Valerie Massardier-Nageotte2,3 Laboratoire des Matériaux Polymériques MultiPhasiques, Faculté de Technologie, Université Ferhat ABBAS, Sétif, Algérie 2 INSA-Lyon, Ingénierie des Matériaux Polymères – IMP, Villeurbanne, France 3 Université de Lyon, Lyon, France

*mounira1990@live.com

Abstract Polylactic acid (PLA) is an attractive candidate for replacing petrochemical polymers because it is from renewable resources. In this study, a specific PLA 2002D was melt-mixed with two plasticizers: triethyl citrate (TEC) and acetyl tributyl citrate (ATBC). The plasticized PLA with various concentrations were analyzed by differential scanning calorimetry (DSC), dynamic mechanical analysis (DMA), melt flow index (MFI), thermogravimetric analysis (TGA), X-ray diffraction (XRD), UV-Visible spectroscopy and plasticizer migration test. Differential scanning calorimetry demonstrated that the addition of TEC and ATBC resulted in a decrease in glass transition temperature (Tg), and the reduction was the largest with the plasticizer having the lowest molecular weight (TEC). Plasticizing effect was also shown by decrease in the dynamic storage modulus and viscosity of plasticized mixtures compared to the treated PLA. The TGA results indicated that ATBC and TEC promoted a decrease in thermal stability of the PLA. The X-ray diffraction showed that the PLA have not polymorphic crystalline transition. Analysis by UV-Visible spectroscopy showed that the two plasticizers: ATBC and TEC have no effect on the color change of the films. The weight loss plasticizer with heating time and at 100°C is lesser than at 135 °C. Migration of TEC and ATBC results in cracks and changed color of material. We have concluded that the higher molecular weight of citrate in the studied exhibited a greater plasticizing effect to the PLA. Keywords: Polylactic acid (PLA), plasticizer, triethyl citrate, acetyl tributyl citrate.

1. Introduction Biobased polyesters have attracted much attention due to their biodegradability and biocompatibility which offer clear advantages for both customers and environment. In recent years, Polylactic acid (PLA) is one of the most promising candidates as partial alternative of petrochemical polymers because it is biodegradable and produced from lactic acid obtained by fermentation of renewable raw materials and readily biodegradable[1]. PLA like most synthetic polymers from petroleum resources, which are rarely used alone by themselves, and needs to improve some properties by an inclusion of additives[2]. Poly (lactic acid) PLA has very low toxicity and high mechanical performance, comparable to other commercial polymers, for example, poly (ethylene terephthalate) (PET) and polystyrene (PS)[3-5]. It is highly transparent and has good barrier properties to aromas. In addition, PLA is also biodegradable, compostable and has good mechanical properties[5,6]. However, its major drawbacks as high cost and intrinsic characteristics (i.e. hard and brittle materials) have hindered its wide spread use, especially in the area of packaging applications[7-9]. Plasticizers are widely used additives for polymeric materials to enhance their flexibility, processability, and ductility. Generally, an efficient plasticizer has to reduce the glass transition temperature (Tg) and melting point of the plasticized materials[10]. Different types of plasticizers have been investigated to improve the flexible properties of PLA such as poly (ethylene glycol) (PEG),

Polímeros, 25(6), 581-590, 2015

citrate esters[11-15], oligomeric lactic acid[16] and triacetine[14,16]. There are several important considerations when choosing a plasticizer for PLA for biomedical applications. It should be a nontoxic substance miscible with the polymer, thus creating a homogeneous blend. Also, the plasticizer should not be prone to migration as this would cause the material to regain the brittleness of pure PLA[17]. Thus, good miscibility between PLA and plasticizers is essential. Various plasticizers, which have been proved as effective plasticizers for PLA, citrate esters are used as plasticizers with a variety of different polymers such as poly (methyl methacrylate)[18] and cellulose acetate[19]. They are nontoxic and approved for use as additives in food, personal care products, and in medical plastics[18]. The reason for good solubility of citrate plasticizers in PLA is due to the polar interactions between the ester groups of PLA and the plasticizer[20]. Two biodegradable, nontoxic plasticizers that have been successfully blended with PLA are triethyl citrate (TEC) and acetyl tributyl citrate (ATBC). At the concentrations used (up to 30%) these esters were shown to be compatible with PLA and to generate significant decrease in glass transition temperature, thereby enhancing the ductility of the material. The aim of this study is to investigate the effects of TEC and ATBC with various molecular weight and contents on thermal, dynamical, rheological and plasticizer migration properties of PLA.

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

Maiza, M., Benaniba, M. T., Quintard, G., & Massardier-Nageotte, V.

2. Materials and Methods

2.3.2 Dynamic mechanical analysis (DMA)

2.1 Materials

Measurements were carried out with a dynamic mechanical analyzer (RSAII) at a frequency of 1 Hz. All films (30 × 7 × 0.5 mm3) were tested at a strain of 0.01% using a 3 °C min–1, temperature ramp from –50 to 120 °C.

Poly (lactic acid) (PLA 2002D, extrusion grade, MFI: 5-7 g 10 min at 210 °C, 2.16 Kg) was provided by Cargill-Dow (USA). The melting temperature (Tm) was 180 °C and the glass transition temperature (Tg) was 60 °C, acetyl tributyl citrate (ATBC) and triethyl citrate (TEC) (Scheme 1a and 1b) were purchased from Sigma Aldrich (France).

2.2 Sample preparation Before processing by melt-blending PLA was dried for 24 h at 60 °C, PLA was blended with the plasticizers in a Brabender plastograph. The temperature was 190 °C and the blending time 7 min, and blend rotation speed was 30 rpm. The plasticizers TEC and ATBC were blended to PLA at 5, 10, 15, 20 and 30% weight. Table 1 summarizes the compositions of each formulation.

2.3 Characterization 2.3.1 Differential scanning calorimetry (DSC) Differential scanning calorimetry (DSC) was conducted on a TA instruments DSC Q10. All samples were exposed to consecutive heating and cooling programs to eliminate their thermal history: first heating from 25 °C to 220 °C, isothermal for 3 min, cooling from 220 °C to –50 °C, isothermal for 3 min, and second heating from –50 °C to 220 °C. The heating and cooling rates were 10 °C min–1.

2.3.3 Melt flow index (MFI) Melt flow index (MFI) is the mass flow rate index, expressed in grams, extruded isothermally 10 min under constant load through a die of standard dimensions was measured by using Melt Flow Indexer (MFI Controlab Melt flow rate apparatas model 5) at 190 °C and 2.16 Kg. 2.3.4 Thermogravimetric analysis (TGA) Thermogravimetric analysis (TGA) of the samples was studied by using TA instrument Q100 (TA). The samples were heated from 30 to 500 °C with the heating rate of 10 °C min–1 under nitrogen atmosphere at the flow rate of 20 ml min–1. 2.3.5 X-ray diffraction (XRD) X-ray diffraction (XRD) of films study was carried out using a Bruker D8 Advance diffractometer with Cu-Ni radiation (λ=1.54184 nm). The diffractogram was scanned in the ranges from 6-70o at a scan rate of 0.05o min–1. 2.3.6 UV-Visible spectroscopy The light transmittance of treated and plasticized PLA with TEC and ATBC were measured with a UV-Visible spectrophotometer instrument (Unicam 300UV; England) in the range of 200-600 nm at room temperature. 2.3.7 Plasticizer migration The samples were placed in a vented oven at isothermal temperatures of 100 and 135 °C. In defined time intervals (200, 400, 600 and 800 min) the samples were weighted. Before weighting the surface were cleaned up from the migrated plasticizer[21]. The weight loss of the plasticizer was assumed as the total weight loss of the samples during heating time (t) and calculated according to Equation 1: = Weight loss (%)

m0 − m t ×100 (1) m0

where m0 is the weight before migration and mt is the weight of after migration at t. Table 1. Composition of treated (Trt) PLA and plasticized PLA with TEC and ATBC plasticizers.

Scheme 1. Chemical structure of: (a) (TEC); (b) (ATBC). 582

Formulation Trt PLA PLA-TEC5 PLA-TEC10 PLA-TEC15 PLA-TEC20 PLA-TEC30 PLA-ATBC5 PLA-ATBC10 PLA-ATBC15 PLA-ATBC20 PLA-ATBC30

PLA (%) 100 95 90 85 80 70 95 90 85 80 70

TEC or ATBC (%) 0 5 10 15 20 30 5 10 15 20 30

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Biobased additive plasticizing Polylactic acid (PLA)

3. Results and Discussion 3.1 Differential scanning calorimetry (DSC) During the last heating scan the glass transition, cold crystallization and melting temperature of the material were determined. The degree of crystallinity (Xc) was calculated from the DSC data using the following relationship: = Xc ( % )

∆H m − ∆H cc ×100 (2) ∆H f ∆X PLA

where ΔHm, ΔHcc and XPLA are the enthalpy of melting, enthalpy of cold crystallization and weight fraction of PLA respectively, ΔHf is the heat of fusion defined as the melting enthalpy of 100% crystalline PLA, which is 93 J.g–1[22]. The differential scanning calorimetry (DSC) results (last scan) for the treated and plasticized PLA in Brabender plastograph are shown in Figure 1 and summarized in Table 2.

The plasticizers decreased the glass transition temperature (Tg) of treated PLA from 60.42 °C to 10.29 °C and to 12.21 °C with addition 30% of TEC and ATBC respectively. As expected, by increasing plasticizer content, a decrease in Tg occurs, which is true for the TEC and ATBC. The low molecular size of the plasticizer allows it to occupy intermolecular spaces between polymer chains, reducing the energy for molecular motion and the formation of hydrogen bonding between the polymer chains, which in turn increases free volume and molecular mobility. By increasing the content of the plasticizer, the effectiveness of the citrate plasticizer to reduce the Tg of the PLA is generally enhanced[20].

The addition of TEC or ATBC to the PLA affects the cold crystallization temperature (Tcc). The Tcc observed in treated PLA at 130.94 °C is depressed to 71.72°C and to 69.26 °C with addition of 30% of TEC and ATBC respectively. The decreasing of Tg, Tcc and melting temperature (Tm) were

Figure 1. DSC diagrams of treated PLA and plasticized PLA with: (a) PLA/TEC; (b) PLA/ATBC at various concentrations. Table 2. Thermal properties and crystallinity of treated and plasticized PLA with TEC and ATBC at various concentrations. Formulation

Tg (°C)

∆Hcc (J/g)

Tcc (°C)

∆Hm (J/g)

Tm (°C)

Xc (%)

Trt PLA

60.42

4.986

130.94

5.240

151.90

0.27

PLA – TEC5

50.48

19.15

120.27

18.46

147.89

0.78

PLA – TEC10

41.20

10.62

114.77

11.46

149.04

1.00

PLA – TEC15

31.21

21.35

104.39

22.48

145.90

1.42

PLA – TEC20

23.56

19.38

91.73

23.07

143.55

4.95

PLA – TEC30

10.29

16.63

71.12

22.22

143.27

8.58

PLA – ATBC5

52.94

22.53

116.03

24.71

146.91

2.46

PLA – ATBC10

44.22

21.33

108.29

24.24

150.04

3.47

PLA – ATBC15

36.67

21.09

101.10

24.73

148.14

4.60

PLA – ATBC20

27.02

20,63

89.03

24.46

146.80

5.14

PLA – ATBC30

12.21

13.97

69.26

19.88

138.41

9.07

Tg: glass transition temperature; Tm: melt temperature; Tcc: cold crystallization; ∆Hm: melting enthalpy; ∆Hcc: enthalpy of the cold crystallization; Xc: crystallinity.

Polímeros, 25(6), 581-590, 2015

583

Maiza, M., Benaniba, M. T., Quintard, G., & Massardier-Nageotte, V. enhanced with a higher plasticizer content as a result from the enhanced chain mobility[23]. The decrease of Tg, Tcc, Tm and the increase of degree of crystallinity as show in Table 2 was seen to depend on the plasticizer content. The percentage of crystallinity of the plasticized PLA is higher than that the treated PLA. For the plasticized PLA with TEC, greatest crystallinity is observed with addition of 30% for a value of 8.58%, and a rate of 9.07% for the plasticized PLA/ATBC. This increase indicates that the crystallization of the PLA becomes easier with the mobility of the chains caused by the citrate ester. The action of a plasticizer is to increase the free volume and to decrease the polymer chain interactions which induce higher chain mobility at lower temperature, the effect linked to plasticization is most probably superposed with a decrease of the glass transition temperature due to chain mobility. Labrecque et al.[19] studied the effect of plasticizers on the thermal properties of PLA and they have found that all plasticizers were miscible with PLA. The reason for good solubility of PLA in citrate plasticizers is due to the polar interactions between the ester groups of PLA and the plasticizer[20].

3.2 Dynamic mechanical analysis (DMA) Curves displaying storage modulus (E’), loss modulus (E”) and loss factor (tan δ) were recorded as a function of temperature for the treated PLA and the plasticized PLA are described in Figures 2, 3 and 4 respectively. Figure 2a and Figure 2b shows the temperature dependence on the storage modulus (E’) for the treated and plasticized PLA with TEC and ATBC respectively. There is a decrease in storage modulus values below Tg as the plasticizer continues to increase the mobility of the polymer chains. The E’ curves for the blends display a short plateau comparing with treated PLA. This indicates a decrease in thermal mechanical stability with additional citrate esters in PLA. The addition of plasticizer in PLA is often reported to

to reduce the stiffness[24]. The thermograms show α-relaxation (main peaks). It is clearly shown in Figure 3a and Figure 3b that the drop in storage modulus following the α-relaxation was found at a lower temperature for the plasticized materials as compared to treated PLA. The curve loss modulus (E”) temperature in Figure 3 provides information about the dispersion and distribution of citrate within PLA phase. Furthermore, a broadening of the width of the peak reﬂecting the glass transition is due to a plasticizer concentration gradient in all the PLA blends, a widening of the width of the curve increased with the plasticizer content (TEC and ATBC). Regarding the decrease of E’ with temperature and the presence of bumps in the tan δ curves after the glass transition, they are due to a cold-crystallization process[25]. We can also note that the peak area of E”, related to the energy required to activate the molecular mobility within the material is different depending on the compositions, the increase of plasticizer content increase the peak area. This suggests that the molecular mobility within the plasticized PLA are easily activated than in the PLA alone. In their work on the phenomena of molecular mobility, David et al.[26], showed that the cooperativity increased with the decrease of disorder and therefore the energy required to activate molecular mobility of a material decreases with the increase of the desorder of one, which is in agreement with the previous result of DSC. The relaxation temperature which can be associated with the glass transition was taken at the maximum of the peak of the damping factor (tan δ). Figure 4 indicates that the Tg was signiﬁcant decreased for all the plasticized PLA. The Tg decreases with the increase of plasticizer content in the PLA. For example, at 30% of TEC and ATBC, the glass transition of PLA decreases from 62.23 °C to 17.49 °C and to 26.91 °C, respectively. These values of glass transition temperature (Tg) obtained by DMA have the same tendency of the results found by the DSC, where Tg is decreased by the addition of plasticizer. Table 3 presents the Tg values for PLA plasticized with

Figure 2. Variation of storage modulus (E’) with temperature of the treated PLA and plasticized PLA with: (a) TEC; (b) ATBC at various concentrations. 584

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Biobased additive plasticizing Polylactic acid (PLA)

Figure 3. Variation of loss modulus (E”) with temperature of the treated PLA and plasticized PLA with: (a) TEC; (b) ATBC at various concentrations.

Figure 4. Variation of loss factor (tan δ) with temperature of the treated PLA and plasticized PLA with: (a) TEC; (b) ATBC at various concentrations.

(TEC or ATBC) obtained by DMA. Furthermore, citrate esters addition also generated broader transition peaks (the full-width half-maximum increased, as indicated in Table 3 and Figure 4). This behavior describes materials with a wide range of relaxation times[27]. The presence of citrate changes the microheterogeneity of the systems, i.e., TEC or ATBC chains can cause the formation of a number of microenvironts with different compositions and different interaction densities in the blends. Typical interactions are hydrogen bonding between the hydroxyl groups from PLA Polímeros, 25(6), 581-590, 2015

and citrate, as well as oxygen, from the citrate esters units, which caused the broadening of the PLA glass transition peak.

3.3 Melt flow index (MFI) The variation of melt flow index (MFI) with the addition of TEC and ATBC was illustrated in Figure 5 respectively. MFI data show that all formulations have higher MFI than of treated PLA (4.6 g per 10 min). The addition of a plasticizer increases the polymer chain mobility which implies a reduction in viscosity and increase of MFI of plasticized PLA[28]. When lower quantity of citrate esters is added in 585

Maiza, M., Benaniba, M. T., Quintard, G., & Massardier-Nageotte, V. Table 3. Glass transition temperature and FWHM of the treated PLA and plasticized PLA with TEC and ATBC determined by DMA. Formulation Trt PLA PLA-TEC5 PLA-TEC10 PLA-TEC15 PLA-TEC20 PLA-TEC30 PLA-ATBC5 PLA-ATBC10 PLA-ATBC15 PLA-ATBC20 PLA-ATBC30

Tg (°C) from tan δ curve 62.23 54.26 45.91 38.71 31.86 17.49 55.24 50.08 40.84 35.69 26.91

FWHM (°C)a 6.48 6.75 7.90 8.70 18.02 20.88 7.02 9.24 10.66 51.18 61.45

Table 4. TGA data of treated and plasticized PLA with TEC and ATBC at various concentrations. Formulation

Ti (°C)

T5% (°C)

T50% (°C)

Tmax (°C)

Trt PLA

295.30

348.94

383.43

390.00

PLA-TEC5

247.51

331.56

383.26

391.43

PLA-TEC10

206.11

303.54

384.52

395.52

PLA-TEC15

171.81

259.43

383.05

394.98

PLA-TEC20

156.53

211.04

377.72

392.25

PLA-TEC30

141.08

193.40

373.98

391.32

PLA-ATBC5

270.07

335.11

385.54

395.52

PLA-ATBC10

214.10

313.63

383.05

393.74

PLA-ATBC15

175.37

260.57

381.30

391.08

PLA-ATBC20

161.86

254.00

378.96

394.27

PLA-ATBC30

146.58

214.36

373.85

390.01

(FWHM): Full-width half-maximum measured from the tan δ curve.

PLA, the small citrate molecule can penetrate PLA matrix and then make PLA matrix slip and flow easier. With the addition of citrate increased large amounts of citrate molecule gathered with each other and surrounded the PLA matrix, which can resist slipping and flowing the PLA molecule chain. When the composition of the plasticizer increases, more molecules of the TEC or ATBC surround the PLA, which becomes more fluid and slippery. The melt index is in good correlation with the results of DSC and DMA.

3.4 Thermogravimetric analysis (TGA) Thermogravimetric analysis (TGA) is an effective approach evaluating the thermal stability of polymeric material. Figure 6a and Figure 6b shows the TG and DTG curves of the neat PLA and plasticized PLA with TEC and ATBC respectively. TGA were carried out on the different formulations reveals that the initial decomposition temperature of plasticized PLA with different concentrations of citrate esters is lower compared to neat PLA. While the kinetics of degradation shows that the plasticized PLA are stable within the range of interest (130 °C). However, the initial decomposition temperature of PLA shifts systematically to lower temperature when citrate esters are added. This shift is globally more important when the amount of plasticizer is higher. For example, a shift from 295 °C to 141 °C and to 146 °C is observed for the initial decomposition temperature of the plasticized PLA with 30% of TEC and ATBC respectively. The initial decomposition temperature (Ti), temperature at 5% and 50% weight loss (T5%) and (T50%) respectively and temperature of maximum decomposition rate (Tmax) of treated and plasticized PLA are shown in Table 4. On increasing the concentration of TEC and ATBC, the T5% and T50% are shifted towards lower values. It was also observed the increased of Tmax of PLA with increased of TEC content, further indicating the role of TEC in promoting the thermal resistance of PLA[29]. The thermal stability of plasticized PLA with ATBC is greather than those with TEC. The Tmax of the plasticized PLA with ATBC is higher than the treated PLA. The presence of hydroxyl end groups in the PLA oligomer chains was found to be critical for degradation, initiating the chain-scission and decreasing 586

Figure 5. Variation of MFI of treated PLA and plasticized PLA with various concentrations of TEC and ATBC.

the thermal stability[30,31]. The range temperature between 284 and 335 °C corresponds to evaporation of plasticizer components, that’s where the boiling point of TEC and ATBC is 127 °C and 173 °C respectively, where they appear clearly in DTG of plasticized PLA with 30% of TEC or ATBC.

3.5 X-ray diffraction (XRD) The X-ray diffraction (XRD) is used to observe changes in the crystallinity of the realized formulations. Figure 7a and Figure 7b show the XRD patterns of treated PLA and plasticized PLA with TEC and ATBC respectively. The PLA has only a strong diffraction at 16.7° assigned to the crystalline phase α, which confirms that the PLA has no polymorphic crystalline transition[32]. The PLA alone and various formulations have the same crystal structure[20]. The diffraction peaks corresponding to the plasticized PLA are offset slightly with increasing TEC and ATBC content and a greater shift for the composition PLA-TEC or ATBC at 30% probably total crystallinity. This indicates that the incorporation of the plasticizer accelerates the Polímeros , 25(6), 581-590, 2015

Biobased additive plasticizing Polylactic acid (PLA)

Figure 6. TGA/DTG thermograms of treated PLA and plasticized PLA with: (a) TEC; (b) ATBC at various concentrations.

Figure 7. X-ray diffractograms for treated PLA and plasticized PLA with: (a) TEC; (b) ATBC at various concentrations.

crystallization of PLA, the same result found by DSC. This can be explained by the finding by Miyajima et al.[33], Li et al.[34], that the chain mobility of amorphous PLA could increase the crystallinity due to the production of PLA with lower molecular weight and shorter chain which was more mobile and more susceptible to crystallize than longer ones.

3.6 UV-Visible spectroscopy To further study the effect of these plasticizers on the transparency of the PLA. The UV-visible light transmittance of treated and plasticized PLA with TEC and ATBC were measured as shows in Figure 8a and Figure 8b. The spectra clearly show that whatever the level of plasticizer, the general appearances of all spectra are identical. All the maximum absorbance range are between about 208‑214 nm. Both plasticizers (TEC and ATBC) have no effect on the color change of the films. In the Polímeros, 25(6), 581-590, 2015

same field, there appear weak staining is due to the effect of film thickness.

3.7 Plasticizer migration The weight loss, assumed as the weight loss of the plasticizer, was calculated as an arithmetic mean of three measured values. The variation of weight loss of treated and plasticized PLA with TEC and ATBC as a function of time at 100 °C and 135 °C is shown in Figure 9a and Figure 9b. The plasticizer loss increases with time and at 100 °C is lower than at 135 °C. It’s clear that higher temperature favors the plasticizer migration. Citrate esters (TEC and ATBC) migrated out of the PLA at elevated temperature. Plasticizer weight loss was directly proportional to the temperature and plasticizer concentration in PLA. Higher temperatures eased plasticizer migration compared to lower temperatures. The rate of weight loss at 135 °C is highest at 400 min for all samples. After 800 min all citrate esters 587

Maiza, M., Benaniba, M. T., Quintard, G., & Massardier-Nageotte, V.

Figure 8. UV- Visible spectra of treated PLA and plasticized PLA with: (a) TEC; (b) ATBC at various concentrations.

Figure 9. Weight loss of treated PLA and plasticized PLA with: (a) PLA/TEC; (b) PLA/ATBC at various concentrations after heating at 100 °C and 135 °C. 588

Polímeros , 25(6), 581-590, 2015

Biobased additive plasticizing Polylactic acid (PLA) migrate from the simple and all the samples show little color changes at 100 °C and 135 °C. The color change may be evidence of the increase of PLA crystallinity and / or thermal degradation. In the case of sample PLA plasticized with TEC is transparent and colorless ATBC becomes opaque and it can be concluded that the change in degree of crystallinity is the main reason. After 800 min at 135 °C the effect of the migration is more visible and the samples are cracked and broken (especially plasticized PLA). The color change of all samples is also more pronounced. In case of loss of plasticizer is worsened flexibility, ie. The samples become more rigid and easy to break.

4. Conclusions In this study, the characterization of plasticized PLA with TEC and ATBC by using the simple melt blending method was reported. The thermal properties of plasticized PLA shows that TEC and ATBC are effective in lowering the glass transition temperature (Tg), the melting temperature (Tm) and the cold crystallization temperature (Tcc) of the PLA. In addition, its crystallinity increases with increasing content of plasticizer. The evaluation of the influence of the type and plasticizer content on the viscoelastic properties of PLA with DMA indicates a decrease in the storage modulus is observed for plasticized PLA, indicating the flexibility and mobility of the amorphous phase of PLA caused by TEC and ATBC. The value of the melt flow index (MFI) was observed for all samples are higher than those of neat PLA. The TGA results indicated that the TEC and ATBC promote a decrease in thermal stability of the PLA. X-ray diffraction shows the appearance of a strong diffraction at 16.7° assigned to the crystalline phase α, which confirms that the PLA has no crystalline polymorph transition. UV‑Visible spectroscopy shows that the two plasticizers: ATBC and TEC have no effect on the color change of the films. Thermally induced migration of citrate esters (TEC and ATBC) from plasticizer PLA was investigated. The weight loss plasticizer with heating time and at 100 °C is lesser than at 135 °C. migration of TEC and ATBC results in cracks and changed color of material., Finally it can be said that the higher molecular weight of citrate exhibited a greater plasticizing effect to the PLA.

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Polímeros , 25(6), 581-590, 2015

http://dx.doi.org/10.1590/0104-1428.2039

Characterization of clay filled poly (butylene terephthalate) nanocomposites prepared by solution blending Khalid Saeed1* and Inayatullah Khan1 1

Department of Chemistry, University of Malakand, Chakdara, Dir (Lower), Khyber Pakhtunkhwa, Pakistan *khalidkhalil2002@yahoo.com

Abstract Kaolin clay/poly (butylene terephthalate) (clay/PBT) composites films were prepared by solution casting technique. The scanning electron microscope (SEM) study showed that clay particles were well dispersed and embedded within the PBT matrix. The TGA thermograms showed that the thermal stability of PBT matrix was slightly improved by the incorporation of clay into the polymer matrix. The polarized optical microscopy (POM) study presented that the size of spherulites of PBT was decreased by the incorporation of clay into matrix, which might be due to nucleation effect of kaolin clay. The tensile strength and modulii of PBT polymer matrix were also significantly improved by the addition of clay polymer matrix. The solvent uptake study showed that the uptake of various solvents by clay/PBT nanocomposite were lower than neat PBT. Keywords: kaolin clay, morphology, nanocomposites.

1. Introduction Clay/Polymer nanocomposites exhibits some superior material properties like improved mechanical and thermal properties, gas permeability and fire retardance as compared to common neat polymers[1-3]. These superior properties of clay/polymer composite are may be due to the large surface area of clay, ionic bond between polymer and clay, and good dispersion of clay in the matrix[4]. Clay/polymer nanocomposites are commonly prepared by three principal methods, like solution intercalation, in-situ polymerization and melt intercalation[5-8]. Melt intercalation is the most convenient, versatile, compatible and environmentally favored technique for the preparation of clay/polymer composite[9,10]. Poly(butylene terephthalate) (PBT) is an important thermoplastic polymer of polyester family. PBT is a semi crystalline engineering thermoplastic with some superior properties like high crystallization rate, chemical and thermal resistance, high impact strength, low molding temperature and excellent processing. Due to these superior properties, PBT has wide application in automotive, electronic, electric and packaging industry[11-14]. Clay/PBT nanocomposites have been studied by many researchers. Chang et al.[15] prepared clay/PBT nanocomposites through In situ interlayer polymerization and found that the dispersion of a very small amount of organoclay can greatly enhanced thermal and mechanical properties of PBT. Xiao et al.[12] prepared clay/PBT nanocomposite via direct melt intercalation method using thermally stable organically modified montmorillonite as filler. They also found the remarkable improvement in melting temperature and rate of crystallization of the resulted clay/PBT nanocomposites. Twin screw extrusion was also used for melt intercalation of clay/PBT nanocomposites[16]. Kaolin was extensively used as filler for plastics and rubber. More than 60% of the word production of koalins is used in the paper industry as fillers for cellulose fibers and as coating particles in paint industry. The coating particles should be smaller than 2 mm and that for paper fibers being

Polímeros, 25(6), 591-595, 2015

5-7 mm. kaolin is largely used in ceramic products like porcelain, bone china, vitreous sanitary ware, earthenware, pipes, tiles and refractory bricks[17,18]. In the present study, kaolin clay was used as filler for PBT polymer matrix. Kaolinite is the principal constituent (85-95%) of kaolin clay. Other clay minerals of kaolin group are nacrite, dickite and are generally represented by molecular formula Al2Si2O5(OH)4[19]. Clay/PBT composite films were prepared by solution casting method in order to get a uniform of clay within the PBT matrix and then studied the effect kaolin clay on crystallization, morphology, mechanical and thermal properties of PBT.

2. Materials and Methods 2.1 Materials PBT (average molecular weight of 38,000) and trifluroacetic acid were purchased from Sigma Aldrich and were used as received. The reinforcement material Kaolin (china clay) was kindly provided by local china clay plant Swat, Pakistan.

2.2 Thermal treatment of kaolin clay Kaolin clay was dehydrated upon heating to 350 °C in the presence of oxygen. Thermal treatment below 400 °C does not lead any structural change in the clay. However, above this temperature dehydroxylation takes place which cause structural changes in kaolin[17]. To avoid these structural changes, clay is heated up to 350 °C, cooled and store for further use.

2.3 Preparation of samples The neat PBT and 0.5, 1, 2 wt% clay/PBT composite samples were prepared via solution casting method. The PBT was first dissolve in trifluoroacetic acid, and then added

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

Saeed, K., & Khan, I. a known quantity of kaolin clay in the polymer solution. The mixture solution was stirred well and then sonicated for about 40 min in order to get homogeneous solution. The smooth dry composite ﬁlms/sheets were obtained after the removal of solvent with distilled water. Similar procedure was used for the preparation of 0.5, 1 and 2 wt % kaolin/PBT composite ﬁlms.

2.4 Characterization The gold coated cryofractured surface of neat PBT and composite samples were analyzed by SEM (model JEOL JSM-5910). The POM study was performed by using POM (modal Optika B-600 POL). The samples were melted on heater, squeeze between two glass slides and then analyzed under POM. The thermal properties were studied by using TG/DTA, Perkin Elmer instrument at heating rate of 20 °C/ min from room temperature to 800 °C under nitrogen atmosphere. The mechanical properties of neat PBT and clay/PBT composites were examined using universal testing machine (UTM), Model 100-500 KN, Iestomeric Inc. The solvent uptake capacity of the samples were studied in selected solvents (distilled water, 2M HNO3 solution, chloroform and kerosene oil). Small dry rectangular strips of the samples were weighed and then immersed in vials containing about 10 mL solvents. Percent swelling is determined by the following equation[20]. = Solvent uptake (%)

Figure 1. SEM micrographs of 2 wt% clay/PBT nanocomposites.

W0 − W ×100 (1) W

Where “W” is the weight of dry sample and “Wo” is the weight of wet sample in solvents.

3. Results and Discussions 3.1 Morphology of clay/PBT nanocomposites The dispersion of fillers within the matrix has greatly affected the properties (such as thermal, mechanical, electrical properties etc.) of the matrixes. Figure 1 illustrates the SEM micrograph of fractured surface (broken in liquid nitrogen) of clay (2 wt %)/PBT composites film. The micrograph presented that the clay particles were not present in agglomerated but dispersed well within the PBT polymer matrix. The Figure 1 also presented that the size of the clay particles were below 800 nm. Figure 2a, b show POM images of neat PBT and 2 wt% clay/PBT nanocomposites, respectively. Thin films of PBT and the composite were prepared by cooling both samples from the melt state. Figure 2a presented almost uniform formation of spherulites throughout the PBT matrix. The shape of neat PBT spherulites is Maltese type. The size of spherulites of neat PBT was below 25 µm. The POM micrograph (Figure 2b) of clay (2 wt %)/PBT nanocomposite showed that the size of spherulites significantly decrease as compare to neat PBT. The decrease in size of spherulites in clay/PBT nanocomposite is might be due to the nucleation effect of kaolin clay. The decrease in size of spherulites in the case of nanocomposites materials were also reported by Saeed and Park, where they incorporated multi walled carbon nanotubes into polycaprolactone[21]. 592

Figure 2. POM images of (a) pure PBT and (b) 2 wt% clay/ PBT.

3.2 Thermal properties of neat PBT and clay/PBT nanocomposites TGA study was performed for the purpose to determine the effect of kaolin on the degradation temperature and thermal stability of PBT. The TGA curves of neat PBT, 0.5, 1.5 and 2 wt% clay/PBT nanocomposites are collectively shown in Figure 3. Figure 3 illustrated that the weight of neat PBT and nanocomposite remain unchanged till 300 °C. After 300 °C, the TG curves starts dipping down from 300 to 450 °C, which shows that the weight loss occurs in this particular range of temperature. The degradation of Polímeros , 25(6), 591-595, 2015

Characterization of clay filled poly (butylene terephthalate) nanocomposites prepared by solution blending neat PBT polymer started at 300 °C and completed at about 430 °C. While in the case of Clay/PBT nanocomposite, the degradation of compsite started at about 350 °C. It was also found that the degradation temperature of clay/PBT samples was slightly shifted to higher temperature (5 to 10 °C). The residual quantities which contribute to the kaolin clay in the clay/PBT composite, remain at higher temperature. Figure 4 illustrates the DTA curves for neat PBT, 0.5, 1.5 and 2 wt % clay/PBT nanocomposite samples. All the samples have a weak endothermic peak around 200 °C and a sharp endothermic peak around 400 °C (Figure 4). The weak endothermic peak around 200 °C (225 °C-227 °C) represents the melting, while the sharp

one around 400 °C may represent the degradation of neat PBT[22]. Multiple peaks in this region are also possible which may be due to recrystalization process[23]. As clear from Figure 4 the most noticeable difference is the height of peaks which increase gradually with increasing clay contents. The strong endothermic peaks in case of 2 wt% clay/PBT nanocomposites are due to the interaction between clay and polymer, which results more stable morphology through heterogeneous nucleation mechanism [24].

3.3 Mechanical properties of neat PBT and clay/PBT nanocomposites The dispersion of fillers plays a key role in improving the mechanical properties of polymeric materials. The well dispersed filler can results superior mechanical properties. Table 1 shows the mechanical properties of neat PBT, 0.5, 1, 1.5 and 2 wt% clay/PBT nanocomposites. The increase in clay content cause a significant increase in tensile strength and modulus values as shown in Table 1. The tensile strength values increase from1.6 N/mm2 (neat PBT) to 5.6 N/mm2 (2 wt %) while tensile modulus increase from 166.6 N/mm2 (neat PBT) to 1080.9 N/mm2 (2 wt %). The results shows that both tensile strength and tensile modulus of PBT was enhanced significantly by incorporation of clay (up to 2 wt %) in to the polymer matrix. Similarly, Gu et al.[25] also reported the improvement of mechanical properties of PET by the addition of modified montmorillonite clay.

3.4 Solvent uptake study of neat nylon 6,6 and clay/nylon 6,6 nanocomposites

Figure 3. TGA curves of neat PBT, 0.5, 1.5 and 2 wt% clay/PBT nanocomposites.

Tables 2-4 illustrate the solvent uptake of neat PBT and clay/PBT nanocomposites. The distilled water, 2 M HNO3 solution, and kerosene were taken as solvents during the solvent uptake study. The results show that initially the Table 1. Mechanical Properties of PBT and clay/PBT nanocomposites. Samples Neat PBT 0.5 wt % clay/PBT 1 wt % clay/PBT 1.5 wt % clay/PBT 2 wt % clay/PBT

Stress Yield (N/mm2) 1.6 ± 0.2 2.2 ± 0.1 3.9 ± 1.26 5.6 ± 0.04 5.6 ± 0.01

Young Modulus (N/mm2) 166.6 ± 4.37 163.6 ± 3.85 443.9 ±7.5 567.6 ± 84.04 1080.9 ± 64.05

Table 2. Percent water uptake by neat PBT and clay/PBT composites.

Figure 4. DTA thermographs of (a) neat PBT, (b) 0.5 wt%, (c) 1.5 wt% and (d) 2 wt% clay/PBT nanocomposites. Polímeros, 25(6), 591-595, 2015

Time (h) ½ 1 2 3 6 12 24 48 72

PBT 100.0 160.0 200.0 240.0 256.0 280.0 320.0 328.0 335.00

0.5 wt% clay/PBT 100.8 150 175 210 225 258.3 275 312 325

1.5 wt% clay/PBT 80 100.3 128 140.7 167.1 200 240 260.1 272.6

2 wt% clay/PBT 33.3 66.7 100.0 133.3 166.6 200.0 200.0 233.3 252

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Saeed, K., & Khan, I. Table 3. Percent kerosene uptake by neat PBT and clay/PBT composites. Time (h) ½ 1 2 3 6 12 24 48 72

PBT 200 220 240 268 300 340 380 420 430

0.5 wt% clay/PBT 166.6 200 233.3 256.2 300.0 333.3 346.6 358.0 361.3

1.5 wt% clay/PBT 150 179 200.8 225 241.6 275.3 305.0 325 332.7

2 wt% clay/ PBT 140.2 160 172.7 200 220.3 235 280 290.9 303

Table 4. Percent HNO3 solution uptake by neat PBT and clay/PBT composites. Time (h) ½ 1 2 3 6 12 24 48 72

PBT 66.6 100.0 133.3 150.0 185.3 200.0 233.7 260.2 266.9

0.5 wt% clay/PBT 50 70.2 100 121.4 150 175.2 200.5 215 222.4

1.5 wt% clay/PBT 50 70 87.4 110 150.9 167.2 190 210 213.7

2 wt% clay/PBT 29.3 53.6 80.00 133.33 156.6 176.0 200.00 200.00 166.6

solvent uptake was rapid and then the uptake of solvents slow down after 24 h. Among the solvents, uptake of distilled water was found more rapid as compare to other kerosene and HNO3 solution. The solvent uptake of neat PBT and clay/PBT nanocomposites is in the following order: Neat PBT > 0.5 wt% clay/PBT > 1 wt% clay/PBT > 1.5 wt% clay/PBT > 2 wt% clay/PBT The Tables 2-4 also presented that solvent uptake decreased as the quantity of kaolin clay increased in PBT polymer matrix. This decrease in the solvent uptake with increase in concentration is might be due to some interaction between the clay and PBT, which leads to cross-linking and tortuosity in the polymer and that inhibit its solvent uptake ability.

4. Conclusions The dispersion of kaolin in the PBT matrix was confirmed by SEM analysis. The neat PBT had maltase shapes spherulites. The size of PBT spherulites was highly reduced by the addition of kaolin clay. The mechanical properties of clay/PBT were significantly enhanced than neat PBT. The thermal stability of PBT was slightly improved (up to 10 °C) by the addition of kaolin into the polymer matrix may also be achieved. It was also found that the solvent uptake by neat PBT than the nanocomposite samples.

5. References 1. Agag, T., & Takeichi, T. (2000). Polybenzoxizine-Montmorillonite Hybrid Nanocomposite synthesis and characterization. 594

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Characterization of clay filled poly (butylene terephthalate) nanocomposites prepared by solution blending Polymer International, 54(2), 348-353. http://dx.doi.org/10.1002/ pi.1686. 17. Jepson, W. B. (1984). Kaolins: their properties and uses. Philosophical Transactions of the Royal Society of London. Series A, Mathematical and Physical Sciences, 311(1517), 411-432. http://dx.doi.org/10.1098/rsta.1984.0037. 18. Murray, H. H. (2000). Traditional and new applications for kaolin, smectite, and palygorskite: a general overview. Applied Clay Science, 17(5-6), 207-221. http://dx.doi.org/10.1016/ S0169-1317(00)00016-8. 19. Grim, R. E. (1968). Clay mineralogy. New York: McGraw-Hill. 20. Haider, S., Park, S. Y., Saeed, K., & Farmer, B. L. (2007). Swelling and electroresponsive characteristics of gelatin immobilized onto multi-walled carbon nanotubes. Sensors and Actuators. B, Chemical, 124(2), 517-528. http://dx.doi. org/10.1016/j.snb.2007.01.024. 21. Saeed, K., & Park, S.-Y. (2007). Preparation and properties of multiwalled carbon nanotube/polycaprolactone nanocomposites. Journal of Applied Polymer Science, 104(3), 1957-1963. http:// dx.doi.org/10.1002/app.25902. 22. Shyang, C. W. (2008). Tensile and thermal properties of poly(Butylene Terephtalate)/organo-montmorillonite. Malaysian

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Polymer Journal, 3(1), 1-13. Retrieved in 16 November 2015, from http://www.cheme.utm.my/mpj/images/080301_1chow. pdf 23. Xiao, J. F., Hu, Y., Wang, Z. Z., Tang, Y., Chen, Z. Y., & Fan, W. C. (2005). Preparation and characterization of poly(butylene terephthalate) nanocomposites from thermally stable organicmodified montmorillonite. European Polymer Journal, 41(5), 1030-1035. http://dx.doi.org/10.1016/j.eurpolymj.2004.11.025. 24. Venkataramani, S., Lee, J. H., Park, M. G., & Kim, S. C. (2009). Structure and properties of polyamide-6 & 6/66 clay nanocomposites. Journal of Macromolecular Science, Part A: Pure and Applied Chemistry, 46(1), 65-73. http://dx.doi. org/10.1080/10601320802515399. 25. Gu, X. H., Zeng, P., Zhou, J. L., & Xu, B. (2014). Preparation and characterization of poly(ethylene terephthalate) incoporated with secondary-modified montmorillonite. Iranain Polymer Journal, 23(4), 249-255. Received: Feb. 03, 2014 Revised: July 27, 2015 Accepted: Aug. 11, 2015

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Síntese e caracterização de MIP com fenilalanina visando sua aplicação na técnica de SPE Synthesis and characterization of MIP with Phenylalanine for their application in SPE Layla Talita de Oliveira Alves1, Cynthia D’Avila Carvalho Erbetta1, Christian Fernandes2, Maria Elisa Scarpelli Ribeiro e Silva1, Roberto Fernando Souza Freitas1 e Ricardo Geraldo Sousa1* Laboratório de Ciência e Tecnologia de Polímeros - LCTP, Escola de Engenharia, Universidade Federal de Minas Gerais - UFMG, Belo Horizonte, MG, Brasil 2 Laboratório de Controle de Qualidade de Medicamentos e Cosméticos, Faculdade de Farmácia, Universidade Federal de Minas Gerais - UFMG, Belo Horizonte, MG, Brasil 1

*sousarg@ufmg.br

Resumo Polímeros Molecularmente Impressos (MIPs) são polímeros sintéticos que apresentam alta seletividade a uma molécula de interesse. O objetivo deste trabalho foi a síntese e caracterização de MIPs para aplicação na extração em fase sólida (SPE), visando a determinação de fenilalanina. Os MIPs foram sintetizados a partir do MAA, fenilalanina, EGDMA, AIBN, em clorofórmio. Também foi sintetizado o polímero não-impresso (NIP), para controle da seletividade dos MIPs. A dessorção da fenilalanina foi realizada em extrator Soxhlet. Os MIPs e NIP foram caracterizados pelas técnicas de análise: FTIR, UV-Vis, MEV, DSC e TG. O MIP apresentou maior capacidade adsortiva à fenilalanina do que o NIP, com uma taxa média de adsorção de 55% comparada a 11% para o NIP. Por MEV o MIP apresentou superfície mais porosa, importante característica para aplicação em SPE. Os estudos realizados mostraram que o MIP sintetizado apresentou grande potencial para aplicação em técnica de SPE. Palavras-chave: síntese e caracterização, MIP, fenilalanina, extração em fase sólida. Abstract Molecularly imprinted polymers (MIPs) are synthetic polymers that have high selectivity to a molecule of interest. The objective of this work was the synthesis and characterization of MIPs for use in solid phase extraction (SPE), in order to determine Phenylalanine. The MIPs were synthesized from the MAA, Phenylalanine, EGDMA AIBN, in chloroform. Non imprinted polymer (NIP) was synthesized to control the selectivity of MIPs. The desorption of Phenylalanine was carried out in Soxhlet extractor. The MIPs and NIP were characterized by the following analytical techniques: FTIR, UV-Vis, SEM, DSC and TG. MIP showed higher adsorption capacity to Phenylalanine than the NIP with an average rate of adsorption of 55% compared to 11% for NIP. SEM MIP showed more porous surface, an important feature for use in SPE. The synthesized MIP in the present study showed great potential for use in SPE technique. Keywords: synthesis and characterization, MIP, phenylalanine, solid phase extraction.

1. Introdução O reconhecimento específico molecular é um requisito fundamental dos sistemas vivos. Assim, não é surpreendente que os cientistas ao longo dos anos tenham investido grandes somas de tempo e esforço na tentativa de mimetizar funções biológicas responsáveis pela seletividade inerente às interações enzima-substrato, antígeno-anticorpo e fármaco-receptor. O conceito de impressão molecular para formação de anticorpos, processo no qual um antígeno era usado como molécula molde (MM) para dar forma à cadeia polipeptídica de anticorpos, surgiu a partir da teoria de Pauling[1]. Em 1949, o trabalho de Dickey põe em prática a teoria de Pauling promovendo a utilização de absorventes específicos com

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memória inerente a uma dada molécula e capazes de se religar seletivamente à mesma[2]. Dessa concepção, idealizou-se produzir um polímero rígido, com seletividade atribuída principalmente a sua estrutura tridimensional complementar à molécula molde, que pudesse atuar de forma similar ao anticorpo, ou seja, que pudesse efetuar seletivamente o reconhecimento molecular[3]. Atualmente, essa técnica encontra-se bem estabelecida, sendo rotineiramente intitulada de Polímeros Molecularmente Impressos ou Molecularly Imprinted Polymers – MIPs, que são polímeros sintéticos que apresentam alta seletividade a uma molécula de interesse. Em geral, são facilmente

Polímeros , 25(6), 596-605, 2015

Síntese e caracterização de MIP com fenilalanina visando sua aplicação na técnica de SPE sintetizáveis, podendo ser moldados de acordo com sua utilização, sendo sua síntese em geral pouco onerosa. Apresentam características vantajosas como estabilidade, robustez, seletividade, resistência a altas temperaturas e pressões e inércia química a ácidos, bases e solventes orgânicos[4]. Em função dessas características, e dada sua alta seletividade, esses polímeros têm sido bastante empregados no preparo de amostras atuando como adsorventes em técnicas de separação, tais como, cromatografia líquida de alto desempenho[5], eletroforese capilar[6], eletrocromatografia capilar[6], cromatografia em camada delgada[7], bem como em extração em fase sólida (SPE)[8] e em microextração em fase sólida[9], sendo ainda muito empregados nas áreas biológica, farmacológica e alimentícia[10]. O emprego de MIPs como materiais adsorventes na técnica de SPE vem adquirindo destaque, sendo esse o principal campo de aplicação desses materiais, pois oferece alto grau de seletividade quando comparado com outros materiais, como cartuchos de sílica modificada (C18) e resinas de troca iônica e, ao mesmo tempo, por serem mais estáveis que os adsorventes de origem biológica[11,12]. A SPE é uma técnica amplamente empregada na extração e concentração de analitos, mesmo quando estão presentes em baixos níveis de concentração, na remoção de compostos interferentes em matrizes complexas e na mudança do meio de solubilização de um analito antes de sua análise cromatográfica. Sendo assim, a SPE é considerada uma técnica muito importante para o preparo de amostras[13]. Quando o material adsorvente é um MIP, a técnica recebe a designação de Molecular Imprinting Solid-Phase Extraction (MISPE)[14]. Vários estudos foram realizados nos últimos anos comprovando a utilidade do método MISPE na extração de compostos em amostras biológicas, biofluidos, tecidos, água, solo e plantas. Os recentes desenvolvimentos em impressão molecular disponibilizaram polímeros que podem ser usados na detecção de fármacos, toxinas, agrotóxicos, aditivos alimentares, compostos tóxicos, entre outros[14]. Um exemplo desses compostos tóxicos para o nosso organismo é a fenilalanina em excesso e seus catabólitos, que têm efeito tóxico nas funções somáticas e do sistema nervoso central, causando anormalidades como falhas no andar ou falar, hiperatividade, tremor, microencefalia, falhas no crescimento e retardo mental, sendo essa doença denominada fenilcetonúria[15]. A fenilcetonúria (PKU) é uma doença hereditária e está relacionada com a falta de uma enzima que é necessária para digerir a fenilalanina. Essa, como não é absorvida, passa a acumular-se no organismo até ser convertida em compostos tóxicos, designados por fenilcetonas (como o fenilacetato e a fenetilamina), que são expelidos pela urina. Os doentes com PKU que ingerem a fenilalanina sofrem de diferentes sintomas de toxicidade, incluindo atrasos mentais (especialmente em crianças) e distúrbios intelectuais nos adultos. A fenilalanina (Phe) é um aminoácido essencial, ou seja, não é sintetizado pelo organismo humano, devendo ser obtido por meio da alimentação. É classificada como um aminoácido apolar e possui um anel aromático em sua cadeia lateral[15,16]. A fenilalanina se apresenta na forma de Polímeros, 25(6), 596-605, 2015

um cristal branco ou pó cristalino, sem odor ou com um leve odor característico e com um sabor levemente amargo. Prontamente solúvel em ácido fórmico, moderadamente solúvel em água e praticamente insolúvel em etanol. Dissolve-se em ácido clorídrico diluído. Apresenta valores de pK de 2,11 e 9,13, ponto isoelétrico (pI) em 5,48 e massa molar igual a 165,18g.mol–1, além de ponto de fusão de 283°C[16]. Os métodos usados para a remoção de fenilalanina baseiam-se na liberação desse aminoácido por hidrólise química ou enzimática, sendo posteriormente removido por tratamentos diferenciados. Vários métodos são utilizados para a remoção da fenilalanina, como adsorção em carvão ativado ou resinas de adsorção, cromatografia de troca iônica, peneira molecular ou filtração em gel, além de desaminação desse aminoácido pela enzima fenilalanina amônia liase. Apesar do uso do carvão ativado comercial ser uma das formas mais usadas para adsorção de aminoácidos, a utilização desse processo gera um aumento considerável no custo do processo, devido à necessidade de regeneração do adsorvente para posterior utilização. Além disso, existe a questão da não especificidade do carvão ativado comercial, que ao ser aplicado na remoção da fenilalanina em matrizes mais complexas, pode remover também outras substâncias, como, por exemplo, outros aminoácidos[15]. A escolha do método deve considerar a praticidade, a reprodutibilidade e a relação custo/eficiência de cada tratamento e ainda apresentar reconstituição e utilização viáveis. Tais limitações justificam a síntese de MIPs para a determinação de fenilalanina. Diante do exposto, o objetivo deste trabalho foi sintetizar e caracterizar polímeros molecularmente impressos e não‑impressos por meio da técnica de polimerização em massa, utilizando como molécula molde nos MIPs a fenilalanina, bem como estudar a aplicação desses polímeros na técnica de extração em fase sólida.

2. Materiais e Métodos 2.1 Síntese dos MIPs e NIP O ácido metacrílico (MAA) foi utilizado como monômero funcional (MF), o etileno glicol dimetacrilato (EGDMA) como agente de reticulação (AR), o 2,2-azobisisobutironitrila (AIBN) como iniciador radicalar (IR), a fenilalanina (Phe) como molécula molde (MM) e o clorofórmio como solvente. A síntese dos MIPs foi realizada utilizando-se MF em excesso com relação à quantidade de Phe, pois, assim, procurou-se garantir que ele pudesse interagir com todos os sítios disponíveis no molde[17]. Optou-se por utilizar o EGDMA devido ao seu emprego em larga escala e com grande sucesso na síntese de MIPs[4,18]. O uso do iniciador AIBN foi devido ao fato dele sofrer homólise sob irradiação, por meio de uma lâmpada de vapor de mercúrio com emissão entre 345-365nm para gerar os radicais livres[17], condição essa utilizada nesse trabalho. O clorofórmio foi empregado como solvente pelo fato de ser apolar e com constante dielétrica baixa, favorecendo, dessa forma, a estabilidade das interações analito-monômero, além de fornecer um meio onde o analito e monômeros fossem solúveis. Nesse estudo variou-se a razão “MM / MF”, de forma a otimizar as condições de síntese, uma vez que esse fator influencia as propriedades de ligação e seletividade dos MIPs. 597

Alves, L. T. O., Erbetta, C. D. C., Fernandes, C., Silva, M. E. S. R., Freitas, R. F. S., & Sousa, R. G. Para o estudo da influência desse fator, foram sintetizados MIPs com diferentes razões “MM / MF”, conforme mostra a Tabela 1. Nessa tabela também são apresentadas as quantidades de reagentes usadas na síntese do NIP-1. Como se pode ver, as quantidades são exatamente iguais às usadas para o MIP-1, exceto pela ausência de fenilalanina. Esse polímero foi usado como controle/comparação na adsorção dessa substância. Todas as sínteses foram realizadas à temperatura de 4°C utilizando banho ultratermostático (Quimis, modelo Q214D2) e irradiação UV com lâmpada de mercúrio (λ=365nm). A Phe, o MAA, o EGDMA e o clorofórmio foram adquiridos da Sigma-Aldrich. O AIBN foi adquirido da Polysciences. O solvente metanol grau HPLC foi adquirido da Merck e o ácido acético proveniente da Synth. Todos foram utilizados como adquiridos, ou seja, sem purificação prévia. A MM foi pesada dentro de um frasco reacional (frasco ampola de 30mL) em balança analítica (Shimadzu, modelo AEL-405M). Em seguida, mediram-se os respectivos volumes de monômero ácido metacrílico e do solvente clorofórmio, por meio de micropipetas (Eppendorf), adicionando-os no frasco reacional contendo a MM e Phe. O frasco foi agitado manualmente onde, posteriormente, foi adicionado o agente reticulante EGDMA e o iniciador radicalar AIBN. Imediatamente após a adição dos reagentes ao frasco reacional, procedeu-se a purga com nitrogênio durante 5 minutos. O frasco foi então selado e colocado em banho ultratermostático (Quimis, modelo Q214D2) a 4ºC adaptado com uma câmara contendo lâmpada UV (365nm/100W). O tempo de polimerização foi de 6 horas[5]. Após a polimerização, o frasco foi retirado do banho e o clorofórmio (sobrenadante) foi retirado com o auxílio de uma pipeta. O polímero formado (sólido branco) foi secado em estufa a vácuo (Vacuoterm 6030A) por 12 horas a 65ºC para remoção de solvente residual. Posteriormente ao processo de secagem, o frasco ampola de síntese foi quebrado e o polímero seco foi submetido à moagem utilizando um moinho analítico (IKA A11BS1). Os polímeros foram peneirados em tamizes, sendo utilizadas as partículas que ficaram na faixa granulométrica entre 0,042 e 0,050mm pelo fato da porosidade dos frits (discos que suportam e limitam o sorvente dentro do cartucho) utilizados em SPE serem normalmente de 20μm[19]. Os MIPs e NIP obtidos foram acondicionados em frascos plásticos fechados e mantidos à temperatura ambiente.

2.2 Caracterização dos MIPs e NIP Os MIPs e NIP foram caracterizados por FTIR, MEV, DSC e TG. Os espectros de absorção molecular na região do infravermelho, na faixa de 4000 a 650cm–1, foram obtidos em espectrômetro FTIR modelo Nicolet 6700 da Thermo

Fisher SCIENTIFIC, usando acessório de reflexão total atenuada (ATR), com 64 varreduras e resolução de 4cm-1 (cristal Ge), à temperatura ambiente (≈20°C). O equipamento de DSC utilizado, um Shimadzu modelo DSC-60, foi programado para aquecimento das amostras da temperatura ambiente (≈20°C) a 300°C, a uma razão de aquecimento de 10°C.min–1. O porta amostra utilizado foi de alumínio selado com furo e o gás de arraste, nitrogênio, com vazão de 50mL.min–1. A massa das amostras analisadas variou entre 5,0 e 10,0mg. A termogravimetria foi realizada com razão de aquecimento de 10ºC.min–1, da temperatura ambiente (≈20°C) até 600ºC, em porta amostra de platina e atmosfera de nitrogênio com vazão de 50mL.min-1. O equipamento empregado foi um Shimadzu modelo TGA-50WS. A massa das amostras analisadas variou entre 5,0 e 10,0mg. As microscopias dos MIPs e NIP foram realizadas em microscópio eletrônico de varredura (Marca: Hitachi, Modelo: TM3000) em 15,0kV. As imagens obtidas foram processadas utilizando o software TM3000. Todas as amostras foram analisadas à temperatura ambiente de 20°C, aproximadamente. Esse microscópio opera sob baixo vácuo, portanto não necessita de metalização das amostras para obtenção das micrografias.

2.3 Extração da fenilalanina dos MIPs Após a síntese dos MIPs, foi necessário proceder à remoção da molécula molde para que pudesse ocorrer a religação da mesma, quando da aplicação final do MIP. O método escolhido nesse trabalho, para a remoção da fenilalanina, foi extração por Soxhlet. Nesse método, sucessivas lavagens com uma mistura de metanol e ácido acético (9:1, v/v) foram utilizadas até a extração completa da molécula molde, monitorada por meio de análise da solução eluente no espectrofotômetro UV-Vis (VARIAN, modelo Cary 50), sendo determinado, dessa forma, o tempo necessário para a extração. Pesaram-se 2,5g do polímero (MIP) resultante da síntese em papel de filtro Whatman, sendo ele fechado em forma de cartucho e transferido para o Soxhlet. Conectou-se ao Soxhlet, um balão tritubulado contendo inicialmente 300mL de solução de metanol e ácido acético 9:1 (v/v). Utilizou-se uma das três vias para coleta do extrato e na outra foi colocado um termômetro para monitoramento da temperatura, programada para 65±2°C utilizando-se manta aquecedora (FANEM, modelo 178). Após 4 horas de extração, foi retirada uma alíquota de 10mL do extrato e a mesma analisada em espectrofotômetro UVVis (VARIAN, modelo Cary 50), em comprimento de onda de 259nm para verificação da presença da fenilalanina. Confirmada a presença da molécula molde, todo o extrato era retirado do balão e substituído por uma nova solução de 300mL de metanol e ácido acético 9:1 (v/v). Ao atingir uma

Tabela 1. Quantidades dos reagentes utilizadas na síntese dos MIPs com fenilalanina e do NIP. MIP / NIP MIP-1 MIP-2 MIP-3 NIP-1

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Fenilalanina

MAA

Razão “MM/MF”

EGDMA

AIBN

Clorofórmio

(g) 0,0172 0,0319 0,0863 0

(µL) 682 682 3410 682

(mg.mL-1) 25,2 46,8 25,3 0

(µL) 4720 4720 23600 4720

(g) 0,0200 0,0200 0,1000 0,0200

(mL) 5 5 25 5

Polímeros , 25(6), 596-605, 2015

Síntese e caracterização de MIP com fenilalanina visando sua aplicação na técnica de SPE concentração em torno de 200μg.mL–1, passou-se a utilizar um volume de 200mL de solução para a extração com o objetivo de reduzir o consumo de solvente. Esse procedimento de extração foi repetido até o extrato não apresentar mais bandas de absorção em leitura no espectrofotômetro UV‑Visível, indicando a ausência da fenilalanina, totalizando, dessa forma, um período de extração de 64 horas. Após a retirada da MM, o cartucho contendo MIP foi lavado com metanol ainda no Soxhlet, para a retirada do excesso de ácido, por um período de 4 horas, e então submetido ao vácuo em estufa à vácuo (Vacuoterm, modelo 6030A) à temperatura de 65±5°C, para remoção do excesso de solvente e secagem.

2.4 Ensaio de adsorção Para a realização deste ensaio, utilizaram-se todos os MIPs e o NIP sintetizados, bem como uma solução‑controle, sem o polímero. Pesaram-se 25, 50 e 100mg, em duplicata, de cada MIP e NIP em Erlenmeyers de 50mL. Na sequência, adicionaram-se, com o auxílio de micropipeta automática Eppendorf, 5mL de solução aquosa de fenilalanina (450μg.mL–1) em cada Erlenmeyer. Os Erlenmeyers foram colocados em banho térmico com agitador orbital (shaker) (NOVA ÉTICA, modelo 304D) com agitação por 72 horas a 25°C. Posteriormente, as amostras contidas nos Erlenmeyers foram transferidas para tubos Falcon de 15mL e centrifugadas em centrífuga (SPINLAB, modelo SL16 RAV) a 4000rpm por 30 minutos. O sobrenadante foi recolhido, filtrado e analisado em espectrofotômetro UV-Visível (VARIAN, modelo Cary 50) a 258nm. Para a realização dessa leitura, preparou-se uma curva analítica para a fenilalanina em solução aquosa, com concentração variando entre 20 e 450μg.mL–1. O mesmo procedimento para a solução-controle, sem polímero, foi realizado.

3. Resultados e Discussão 3.1 Síntese dos MIPs e NIP Os MIPs e o NIP apresentaram aspecto cristalino, esbranquiçado e estrutura rígida. A formação do possível complexo “monômero funcional - molécula molde” (“MF-MM”), por meio de impressão não-covalente, é ilustrada

na Figura 1. Não foi possível realizar a análise estrutural dos sítios de ligação dos MIPs devido à natureza amorfa do material e à distribuição heterogênea das estruturas dos sítios de ligações.

3.2 Caracterização dos MIPs e NIP Os espectros FTIR referentes à fenilalanina e aos polímeros NIP-1 e MIP-1, MIP-2 e MIP-3, antes da extração da Phe, são apresentados na Figura 2. As principais bandas de absorção características dos grupos químicos presentes nos polímeros e na fenilalanina são apresentadas na Tabela 2. Nos espectros FTIR de todos os polímeros sintetizados e apresentados na Figura 2, foi possível notar as bandas características dos grupos químicos presentes nesses polímeros, com destaque para as bandas entre 2964 e 2957cm–1, característica da deformação axial das ligações de C-H de alifático, em 1722cm–1, característica do estiramento da carbonila (C=O) de ácido carboxílico, e entre 1162 e 1153cm–1, característica da deformação axial de C-O de ácido carboxílico. Nos espectros FTIR para os MIPs e, diferentemente no espectro FTIR para o NIP, observam-se as bandas de absorção características de aminoácidos em 2414 e 2124cm–1, deformação axial das ligações de C-H de anel aromático em 3067cm-1 e a banda de absorção características da ligação C=O de aminoácidos em 1563cm–1. Também pode-se observar que essas bandas apresentaram-se mais intensas nos espectros FTIR para os MIPs 1 e 3, apresentando-se em menor intensidade para o MIP-2, o que permite inferir que a quantidade da Phe nesse MIP pode ser menor. As micrografias obtidas para o NIP-1 e os MIP-1, MIP-2 e MIP-3, antes da extração da Phe, são apresentadas na Figura 3. Pode-se observar nessas micrografias que as polimerizações efetuadas levaram à produção de partículas irregulares. Tal morfologia é característica de MIPs e NIPs sintetizados pelo método de polimerização em massa. Aparentemente não se observou nenhuma diferença entre a micrografia para o NIP-1 e para os MIP-1, MIP-2 e MIP‑3. A presença da Phe parece não ter afetado a morfologia superficial do polímero.

Figura 1. Representação da provável formação do complexo “MF-MM”. Polímeros, 25(6), 596-605, 2015

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Alves, L. T. O., Erbetta, C. D. C., Fernandes, C., Silva, M. E. S. R., Freitas, R. F. S., & Sousa, R. G.

Figura 2. Espectros FTIR para a Phe, o NIP-1 e os MIP-1, MIP-2 e MIP-3 antes da extração da Phe.

Figura 3. Micrografias do NIP-1 (a), MIP-1 (b), MIP-2 (c) e MIP-3 (d), antes da extração da Phe. Tabela 2. Atribuições das principais bandas de absorção no espectro FTIR para os polímeros e a Phe. Número de onda observado (cm–1) Phe MIPs e NIP 849, 778, 746, 699 757 950-944 1074 1155-148 1307, 1163 1300-1250 1453 1563 1625, 1500 1638 1720 2124 2414 2900 2964 2960-2952 3067 -

600

Atribuição deformação angular de C–H de anel aromático proveniente da ligação C-Cl deformação angular de O-H de ácido carboxílico estiramento de C-N de amina primária deformação axial de C-O de éster deformação axial de C-O de ácido carboxílico deformação axial C-C de anel aromático estiramento de C=O de aminoácidos deformação axial de compostos aromáticos estiramento dos grupos vinila C=C estiramento da carbonila (C=O) de ácido carboxílico deformação axial de aminoácidos NH3+ deformação axial de aminas com cargas positivas deformação axial de O-H de ácido carboxílico. deformação axial das ligações de C-H de alifático deformação axial das ligações de C-H de anel aromático

Polímeros , 25(6), 596-605, 2015

Síntese e caracterização de MIP com fenilalanina visando sua aplicação na técnica de SPE Observando as curvas TG para os MIPs (antes da extração) e o NIP-1 sintetizados (Figura 4), percebe-se que o perfil das curvas TG para todas essas amostras é bem semelhante, apresentando uma perda de massa entre a temperatura ambiente e 160ºC e duas perdas de massa sequenciais entre 200 e 500ºC. A primeira perda de massa, em torno de 22±3,4% (exceto para o MIP-3), pode estar relacionada à presença de agente reticulante não reagido, visto que seu ponto de ebulição encontra-se entre 98 e 100°C, e/ou algum solvente residual. Tal fato justifica uma maior perda, em torno de 40%, para o MIP-3, que teve a quantidade de EGDMA aumentada em cinco vezes em relação aos outros MIPs e NIP. As outras duas perdas de massa, que totalizaram cerca de 75±2,4% e 60%, aproximadamente, para o MIP-3, devem estar relacionadas à degradação das cadeias poliméricas dos MIPs e NIP e, também, da fenilalanina, no caso dos MIPs. As curvas DSC para a Phe, o NIP-1 e os MIP-1, MIP-2 e MIP-3 antes da extração da Phe, são apresentadas na Figura 5. Pode-se notar na curva DSC para a molécula molde (Phe), apresentada na Figura 5, existência de 3 eventos endotérmicos muito próximos entre 200 e 300ºC. Esses eventos podem estar relacionados à fusão da fenilalanina e de sua degradação, observada na curva TG por meio das duas perdas de massa

consecutivas. Esses dois eventos térmicos para Phe, fusão e degradação, parece que ocorrem em temperaturas muito próximas. As curvas DSC dos MIPs antes da extração da molécula molde e do NIP sintetizados, (Figura 5), apresentaram basicamente o mesmo perfil. Percebe-se claramente um evento exotérmico entre 80 e 160ºC. Acredita-se que esse evento possa estar relacionado ao fluxo líquido de calor devido à polimerização de moléculas de ácido metacrílico, ainda presentes na estrutura dos MIPs e NIP, e à vaporização do EGDMA e/ou algum solvente residual, ambos fenômenos ocorrendo com o aumento da temperatura e dentro do mesmo intervalo. A polimerização do MAA é extremamente exotérmica, prevalecendo sobre o evento endotérmico envolvendo a perda de massa, conforme foi observado nas curvas TG desses polímeros. As curvas DSC para os MIPs não apresentaram picos relacionados à fusão e degradação da Phe presente nesses polímeros. Talvez isso tenha ocorrido devido às interações dessa substância com as cadeias poliméricas dos MIPs, deixando de apresentar uma fusão cristalina e sua degradação ocorrendo em temperatura superiores, concomitante com as cadeias dos polímeros. A partir das curvas DSC (Figura 5), não foi possível observar a temperatura de transição vítrea desses polímeros. Essa transição, caso exista, seria melhor visualizada em uma 2ª corrida desses polímeros por DSC, análise que não foi executada nesse trabalho, pois isso apagaria a história térmica dos polímeros, dificultando o estudo da influência do processo de síntese realizado no comportamento deles.

3.3 Extração da fenilalanina dos MIPs Os espectros FTIR referentes à fenilalanina e aos MIP-1, MIP-2 e MIP-3, após a extração da Phe, são apresentados na Figura 6.

Figura 4. Curvas TG para a Phe, o NIP-1 e os MIP-1, MIP-2 e MIP-3 antes da extração da Phe.

As bandas características dos monômeros utilizados na síntese e que podem ser visualizadas nos espectros FTIR dos polímeros após a extração, apresentados na Figura 6, são as de EGDMA e MAA na região de 1700cm–1, correspondentes ao estiramento de C=O do grupo éster e carboxila do ácido, em 1150cm-1, correspondente à deformação axial de C-O de éster e na região de 1300 a 1200cm–1, característica de deformação axial de C-O de ácido carboxílico. Na região de 2900cm–1 foram evidenciadas bandas de estiramento de O-H de ácido carboxílico e C-H. Além disso, foi possível observar nos espectros FTIR para os MIPs, após a extração (Figura 6), a ausência das bandas características da Phe como, por exemplo, em 2414 e 2124cm–1, que são bandas características de aminoácidos, além da banda em 3067cm–1, correspondente ao estiramento das ligações de C-H de anel aromático, o que é um forte indício da extração da fenilalanina. As micrografias obtidas para o MIP-1, antes e após a extração da Phe, são apresentadas na Figura 7. Nota-se que o processo de extração de Phe afetou a morfologia do polímero, deixando-o com o aspecto mais poroso, após a extração.

Figura 5. Curvas DSC para a Phe, o NIP-1 e os MIP-1, MIP-2 e MIP-3 antes da extração da Phe. Polímeros, 25(6), 596-605, 2015

As curvas TG para os MIPs 1, 2 e 3, após a extração, bem como para a fenilalanina, são apresentadas na Figura 8. Nota-se nas curvas TG para os MIPs o mesmo perfil de perda de massa, sendo possível observar três eventos de perda de 601

Alves, L. T. O., Erbetta, C. D. C., Fernandes, C., Silva, M. E. S. R., Freitas, R. F. S., & Sousa, R. G.

Figura 6. Espectros FTIR para a Phe e os MIP-1, MIP-2 e MIP-3 após extração da fenilalanina.

Figura 7. Micrografias do MIP-1 antes da extração da Phe (a) e após a extração (b).

massa. No primeiro evento, da temperatura ambiente até 200°C, ocorre uma perda de massa entre 1 e 3% da quantidade inicial de massa analisada, que pode estar relacionada com a presença de solvente residual. Entre 200 e 500°C observam-se duas perdas de massa consecutivas em torno de 95%, possivelmente proveniente da decomposição de metacrilatos (degradação das cadeias poliméricas dos MIPs). Comparando as curvas TG dos MIPs, após extração da Phe, com a curva TG da Phe, percebe-se que esses polímeros não contêm essa substância, indicando que o processo de extração dela foi eficiente. 602

As curvas DSC para a Phe e os MIP-1, MIP-2 e MIP-3 após a extração da Phe, são apresentadas na Figura 9. Percebe-se um evento endotérmico entre a temperatura ambiente e 120ºC nas curvas DSC para os MIPs, possivelmente devido à presença de solvente residual proveniente do processo de extração da fenilalanina. A partir de 180ºC nota-se o início de outro evento endotérmico, pouco intenso e largo, talvez relacionado ao início da fusão e degradação dos polímeros. Comparando a curva DSC para a Phe com as curvas DSC para os MIPs, poderia utilizar-se da suposição Polímeros , 25(6), 596-605, 2015

Síntese e caracterização de MIP com fenilalanina visando sua aplicação na técnica de SPE feita anteriormente, para a situação dos MIPs antes da extração da Phe. No entanto, em função dos resultados apresentados por FTIR e TG, acredita-se que ela não esteja presente nesses polímeros.

3.4 Ensaio de adsorção

Figura 8. Curvas TG para a Phe e os MIP-1, MIP-2 e MIP-3 após a extração da Phe.

O ensaio de adsorção foi realizado de forma a avaliar a capacidade de reconhecimento molecular dos polímeros sintetizados. Assim, a adsorção da fenilalanina nos MIPs (após a extração de Phe) e NIP foi verificada variando a quantidade de polímero e mantendo a concentração inicial da solução entre 448,9 e 449,8μg.mL–1. Os cálculos para o ensaio de adsorção foram feitos determinando a concentração livre de Phe no sobrenadante, de acordo com a curva analítica de Phe em água construída antes do ensaio. A concentração de Phe adsorvida (C, em μg.mL–1) pelos MIPs e NIP foi determinada pela diferença entre a concentração inicial da solução (I, em μg.mL–1) e a concentração de analito livre (F, em μg.mL–1). A taxa de adsorção da Phe (Tad, em %) foi calculada pela relação: Tad = (C / I) x 100. Os resultados do estudo de adsorção contendo a taxa de adsorção da fenilalanina nos polímeros MIPs e NIP são apresentados nas Tabelas de 3 a 6. As taxas de adsorção da fenilalanina em função das massas dos polímeros MIP-1, MIP-2, MIP-3 e NIP-1 utilizadas durante o ensaio, são apresentadas na Figura 10.

Figura 9. Curvas DSC para a Phe e os MIP-1, MIP-2 e MIP-3 após a extração da Phe.

Figura 10. Taxas de adsorção da fenilalanina em função das massas dos MIP-1, MIP-2, MIP-3 e NIP-1.

Tabela 3. Ensaio de adsorção variando a quantidade (massa) do polímero MIP-1. Massa

(I)

(F)

(C)

Tad

(mg)

(μg.mL–1)

(%)

Média e desvio padrão da Tad(%)

25 25 50 50 100 100

448,9 448,9 448,9 448,9 448,9 448,9

411,4 423,3 267,2 284,3 207,5 194,9

37,5 25,6 181,7 164,6 241,4 254

8,35 5,70 40,48 36,67 53,78 56,58

7,03 ± 1,9 38,58 ± 2,7 55,18 ± 2,0

Média e desvio padrão da Tad(%)

Tabela 4. Ensaio de adsorção variando a quantidade (massa) do polímero MIP-2. Massa

(I)

(F)

(C)

Tad

(mg) 25 25 50 50 100 100

(μg.mL–1) 449,8 449,8 449,8 449,8 449,8 449,8

(μg.mL–1) 446,9 448,9 446,8 447,4 444,1 448,4

(μg.mL–1) 2,9 0,9 3 2,4 5,7 1,4

(%) 0,64 0,20 0,67 0,53 1,27 0,31

Polímeros, 25(6), 596-605, 2015

0,42 ± 0,3 0,60 ± 0,1 0,79 ± 0,7

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Alves, L. T. O., Erbetta, C. D. C., Fernandes, C., Silva, M. E. S. R., Freitas, R. F. S., & Sousa, R. G. Tabela 5. Ensaio de adsorção variando a quantidade (massa) do polímero MIP-3. Massa

(I)

(F)

(C)

Tad

(mg) 25 25 50 50 100 100

(μg.mL–1) 449,8 449,8 449,8 449,8 449,8 449,8

(μg.mL–1) 401,3 399,1 348,1 339,8 309,6 312,9

(μg.mL–1) 48,5 50,7 101,7 110 140,2 136,9

(%) 10,78 11,27 22,61 24,46 31,17 30,4

Média e desvio padrão da Tad(%) 11,03 ± 0,3 23,54 ± 1,3 30,81 ± 0,5

Tabela 6. Ensaio de adsorção variando a quantidade (massa) do polímero NIP-1. Massa

(I)

(F)

(C)

Tad

(mg) 25 25 50 50 100 100

(μg.mL–1) 449,8 449,8 449,8 449,8 449,8 449,8

(μg.mL–1) 416,3 419,5 403,6 401,3 401,8 400,4

(μg.mL–1) 33,5 30,3 46,2 48,5 48,0 49,4

(%) 7,45 6,74 10,27 10,78 10,67 10,98

A partir das Tabelas 3 a 6 ou da Figura 10, percebe-se que, para os MIPs 1 e 3, com o aumento da massa de polímero de 25 para 100mg, a quantidade de fenilalanina adsorvida também aumenta, passando de 7,0±1,9 para 55,2±2,0 e de 11,0±0,3 para 30,8±0,5, respectivamente. Esses resultados indicam que, mantendo-se fixa a relação entre a quantidade de MM e MF, o aumento de massa de polímero nesses MIPs tende a aumentar a quantidade de Phe adsorvida de forma bem significativa, dentro da variação de massa estudada. Comparando-se os MIPs 1 e 3, nota-se que as quantidades de MM, MF, AR, IR e solvente (Tabela 1) foram aumentadas em 5 vezes do MIP-1 para o MIP-3, sendo mantido constante o tempo de reação em 6 horas. Ambos apresentaram a mesma tendência quanto à adsorção de fenilalanina, porém o MIP-1 adsorveu uma quantidade maior de Phe com o aumento de massa do polímero, exceto para 25mg (7,0±1,9 para o MIP-1 e 11,0±0,3 para o MIP-3). Apesar da relação estequiométrica ter sido mantida constante para todas essas variáveis de síntese, em termos de estrutura obtida durante a polimerização, parece que foram obtidos mais sítios específicos no MIP-1, para permitir a interação com a Phe. Esse maior número de sítios pode ser devido à reação de formação desses polímeros ser um processo aleatório, permitindo essa diferença. Ou então, pode ser que o tempo de reação para o MIP-3 tenha sido menor do que o necessário para a formação de número de sítios similar ou próximo dos obtidos para o MIP-1. Para melhor entendimento dessa questão é necessária a obtenção de outros MIPs variando-se as condições de síntese, bem como outros ensaios de adsorção de Phe. Ao comparar os resultados de taxa de adsorção de Phe obtidos com os MIPs 1 e 3 com o NIP 1 (apresentados nas Tabelas 3, 5, e 6 e na Figura 10), percebe-se claramente a influência da fenilalanina na formação de sítios específicos e sua interação com o MF durante o processo de polimerização dos MIPs. A única diferença entre os MIPs 1 e 3 com o NIP 1 é que esse último foi sintetizado sem a Phe. As quantidades usadas nas variáveis de síntese (Tabela 1) foram as mesmas, exceto para o MIP-2 que teve um aumento na massa de Phe de 1,85 vezes em relação ao MIP-1 e para o MIP-3, que teve todos os valores dessas variáveis multiplicados por 5. 604

Média e desvio padrão da Tad(%) 7,1 ± 0,5 10,53 ± 0,4 10,83 ± 0,2

Os MIPs, após extração da Phe, apresentaram a mesma taxa de adsorção, aproximadamente, com 25mg de massa de polímero. Porém, com o aumento dessa massa para 50 e 100mg a taxa de adsorção de Phe aumentou significativamente. Para o MIP-1, comparado ao seu polímero controle, NIP-1, o aumento foi de 28% e 44%, respectivamente. Já o MIP-3 apresentou um acréscimo de 13% e 20%, respectivamente, em relação ao seu NIP controle, NIP-1. Portanto, há evidências de que o MF (ácido metacrílico) sintetizado com a MM (fenilalanina) permite um maior número de interações específicas MF-MM no polímero sintetizado com a MM e após extração da mesma desse polímero do que o polímero sintetizado sem a MM. Dentre os MIPs sintetizados, o MIP-2 foi o polímero impresso que menos adsorveu a MM, observando-se taxas de adsorção abaixo de 2%. Quando comparado com a taxa de adsorção observada para o NIP-1, percebe-se que o MIP-2 não foi seletivo para fenilalanina, visto que a taxa mínima de adsorção do NIP foi de 7,1±0,5%. Tal fato pode ser explicado em virtude da variável molécula molde, no MIP-2, ter sido aumentada quase o dobro (1,85) em relação ao MIP-1, enquanto que a quantidade de monômero funcional não foi alterada. Para que se obtenha MIPs com sítios mais específicos, é necessário que se tenha uma quantidade suficiente de MF para complexar com todos os possíveis sítios ativos da MM e, assim, obter sítios específicos. A interação analito-monômero é governada por um processo em equilíbrio, quantidades superiores do monômero em relação ao analito devem ser empregadas com intuito de deslocar o equilíbrio, formando assim maior quantidade de complexos “analito-monômero”. Nesse caso, o aumento da MM pode ter desfavorecido a formação desse complexo, impedindo ou diminuindo drasticamente a criação dos sítios específicos para essa interação. A partir desses resultados, pode-se inferir que uma maior massa de polímero proporciona maior área de contato, e, portanto, maior adsorção da molécula durante um mesmo intervalo de tempo. Polímeros , 25(6), 596-605, 2015

Síntese e caracterização de MIP com fenilalanina visando sua aplicação na técnica de SPE

4. Conclusão Os MIPs e NIP foram sintetizados e caracterizados por FTIR, MEV, TG e DSC. Antes da extração da Phe foi possível observar, por FTIR, as bandas características de seus grupos químicos nos MIPs, com destaque para as bandas de aminoácidos. O processo de extração da Phe dos MIPs foi eficiente, uma vez que os espectros FTIR e as curvas TG, obtidos para os MIPs após a extração, não indicaram a presença da Phe. Conforme observado por MEV, um dos MIPs sintetizado teve sua morfologia superficial afetada pelo processo de extração da Phe, deixando-o com o aspecto mais poroso. Quanto ao estudo de adsorção, os MIPs apresentaram maior capacidade adsortiva pela Phe, obtendo-se taxa média de adsorção de até 56%. Comparando-os ao NIP, para o qual foi obtida taxa de adsorção próxima de 11%, foi possível concluir que esse MIP apresentou-se específico para a determinação da MM Phe, com grande potencial para aplicação como adsorvente em SPE.

5. Agradecimentos Os autores agradecem à FAPEMIG, CAPES, CNPq, CEMIG e PETROBRAS pelo suporte financeiro.

6. Referências 1. Pauling, L. J. A. (1940). Theory of the structure and process of formation of antibodies. Journal of the American Chemical Society, 62(10), 2643-2657. http://dx.doi.org/10.1021/ ja01867a018. 2. Dickey, F. H. (1949). The preparation of specific adsorbents. Proceedings of the National Academy of Sciences of the United States of America, 35(5), 227-229. http://dx.doi.org/10.1073/ pnas.35.5.227. PMid:16578311. 3. Fernandes, R. M. T. (2012). Polímeros de impressão molecular para extração seletiva de fármacos em matrizes biológicas e determinação por LC-MS/MS e MS/MS (Tese de doutorado). Universidade Estadual de Campinas, Campinas. 4. Tarley, C. R. T., Sotomayor, M. D. P. T., & Kubota, L. T. (2005). Polímeros biomiméticos em química analítica. Parte 1: Preparo e aplicações de MIP (“Molecularly Imprinted Polymers”) em técnicas de extração e separação. Quimica Nova, 28(6), 10761086. http://dx.doi.org/10.1590/S0100-40422005000600024. 5. Hung, C., Huang, H., & Hwang, C. (2005). Chiral separations of mandelic acid by HPLC using molecularly imprinted polymers. Eclética Química, 30, 67-73. http://dx.doi.org/10.1590/S010046702005000400009. 6. Suedee, R., Seechamnanturakit, V., Canyuk, B., Ovatlarnporn, C., & Martin, G. P. (2006). Temperature sensitive dopamine-imprinted (N,N-methylene-bis-acrylamide cross-linked) polymer and its potential application to the selective extraction of adrenergic drugs from urine. Journal of Chromatography. A, 1114(2), 239-249. http://dx.doi.org/10.1016/j.chroma.2006.02.033. PMid:16530207.

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7. Mosbach, K. (1994). Molecular imprinting. Trends in Biochemical Sciences, 19(1), 9-14. http://dx.doi.org/10.1016/09680004(94)90166-X. PMid:8140624. 8. Andersson, L., Paprica, A., & Arvidsson, T. (1997). A highly selective solid phase extraction sorbent for pre-concentration of sameridine made by molecular imprinting. Chromatographia, 46(1), 57-62. http://dx.doi.org/10.1007%2FBF02490930. 9. Koster, E. H. M., Crescenzi, C., den Hoedt, W., Ensing, K., & de Jong, G. J. (2001). Fibers coated with molecularly imprinted polymers for solid-phase microextraction. Analytical Chemistry, 73(13), 3140-3145. http://dx.doi.org/10.1021/ac001331x. PMid:11467565. 10. Martín-Esteban, A. (2001). Molecularly imprinted polymers: new molecular recognition materials for selective solid-phase extraction of organic compounds. Fresenius’ Journal of Analytical Chemistry, 370(7), 795-802. http://dx.doi.org/10.1007/ s002160100854. PMid:11569855. 11. Caro, E., Marce, R. M., Borrull, F., Cormack, P., & Sherrington, D. (2006). Application of molecurlay imprinted polymers to solid phase extraction of compounds from environmental and biological samples. Trends in Analytical Chemistry, 25(2), 143-153. http://dx.doi.org/10.1016/j.trac.2005.05.008. 12. Tamayo, F. G., Turiel, E., & Martín-Esteban, A. (2007). Molecularly imprinted polymers for solid-phase extraction and solid-phase microextraction: Recent developments and future trends. Journal of Chromatography. A, 1152(1-2), 32-40. http:// dx.doi.org/10.1016/j.chroma.2006.08.095. PMid:17010356. 13. Yang, J., Hu, Y., Cai, J. B., Zhu, X. L., Su, Q. D. (2006). A new molecularly imprinted polymer for selective extraction of cotinine from urine by solid-phase extraction. Analytical and Bioanalytical Chemistry, 384(3), 761-768. http://dx.doi. org/10.1007%2Fs00216-005-0221-4. 14. Pérez-Moral, N., & Mayes, A. G. (2004). Comparative study of imprinted polymer particles prepared by different polymerisation methods. Analytica Chimica Acta, 504(1), 15-21. http://dx.doi. org/10.1016/S0003-2670(03)00533-6. 15. Clark, H. L. M. (2010). Remoção de fenilalanina por adsorvente produzido a partir da torta prensada de grãos defeituosos de café (Dissertação de mestrado). Universidade Federal de Minas Gerais, Belo Horizonte. 16. Nelson, D. L., & Cox, M. M. (2011). Princípios de bioquímica de Lehninger (5. ed.). Porto Alegre: Artmed. 17. Spivak, D. A. (2005). Optimization, evaluation, and characterization of molecularly imprinted polymers. Advanced Drug Delivery Reviews, 57(12), 1779-1794. http://dx.doi.org/10.1016/j. addr.2005.07.012. PMid:16260064. 18. Al-Kindy, S., Badía, R., Suárez-Rodríguez, J. L., & Díaz-García, M. E. (2000). Molecularly imprinted polymers and optical sensing applications. Critical Reviews in Analytical Chemistry, 30(4), 291-309. http://dx.doi.org/10.1080/10408340008984162. 19. Peçanha, B. R. B., Dias, L. R. S., Spinelli, E., & Muri, E. M. F. (2013). Polímeros de impressão molecular obtidos através de polimerização por precipitação e sua aplicação na técnica de extração em fase sólida. Polímeros: Ciência e Tecnologia, 23(4), 509-513. http://dx.doi.org/10.4322/polimeros.2013.055. Enviado: Fev. 17, 2015 Revisado: Maio 15, 2015 Aceito: Jun. 01, 2015

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http://dx.doi.org/10.1590/0104-1428.1967

T T T T T T T T T T T T T T T T T T

Incorporação de fonte de nitrogênio em partículas de PVA e alginato de sódio e estudo da influência de ciclos de congelamento/descongelamento na caracterização do produto A source of nitrogen incorporation in particulate PVA and sodium alginate and study of the influence of freezing/thawing cycles in the characterization of the product Sinara Queli Welter Nardi1*, Sirlei Dias Teixeira1 e Cristiane Regina Budziak Parabocz1 Departamento de Química, Universidade Tecnológica Federal do Paraná – UTFPR, Pato Branco, PR, Brasil

*welter.sinara@gmail.com

Resumo Neste trabalho foram incorporadas duas fontes de nitrogênio (ureia e caulinita intercalada com ureia) em matriz polimérica de álcool polivinílico e alginato de sódio na proporção de 3:1, utilizando a metodologia de gotejamento em solução de CaCl2. As partículas foram submetidas ao congelamento e posterior descongelamento com o intuito de melhorar a estrutura e resistência térmica da matriz polimérica. As partículas foram caracterizadas através de Análise Elementar, FTIR, DRX e Análise Térmica. As partículas que apresentaram as melhores formulações foram as de álcool polivinílico+alginato de sódio+ureia, pois apresentaram eficiência de incorporação próximas as das partículas de álcool polivinílico+alginato de sódio+caulinita intercalada, mas com maior estabilidade térmica, cerca de 200 °C. Palavras-chave: álcool polivinílico, alginato de sódio, ciclos de congelamento/descongelamento, incorporação de nitrogênio, ureia. Abstract In this work two nitrogen sources (urea and urea intercalated kaolinite) were incorporated in the polymer matrix polyvinyl alcohol and sodium alginate in the ratio of 3: 1, using the methodology drip in CaCl2. The particles were subjected to freezing and subsequent thawing in order to improve the structure and thermal resistance of the polymer matrix. The particles were characterized by elemental analysis, FTIR, XRD, and thermal analysis. The particles that showed the best formulations were those of polyvinyl alcohol+sodium alginate+urea, because they showed the efficiency of incorporation nearby particles polyvinyl alcohol+sodium alginate+urea intercalated kaolinite, but with greater thermal stability, about 200 °C. Keywords: polyvinyl alcohol, sodium alginate, cycles of freeze/thawing, incorporation of nitrogen, urea.

1. Introdução A utilização de polímeros para encapsulamento de diversos materiais tem sido assunto de grande interesse, nas últimas décadas[1,2]. Dois polímeros amplamente utilizados para a incorporação de materiais são o álcool polivinílico (PVA) e o alginato de sódio (AS), por apresentarem baixa toxicidade e alta biodegradabilidade[1]. O PVA é um polímero sintético utilizado em uma ampla gama de aplicações industriais, comerciais, médicas e de alimentos[3]. É um polímero de baixo custo e boa estabilidade térmica, tendo o início de sua

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decomposição entre 230-290 °C[2,4,5]. O AS é um polímero linear polianiônico encontrado em algas castanhas[1]. Sua estrutura consiste em cadeias lineares de monômeros de ácido β-D-manurônico e de ácido α-L-gulurônico, unidos por ligações tipo (1→4) em várias proporções[2,6]. Sua estabilidade térmica é de 210-240 °C[4,5,7,8]. O AS, pode sofrer o processo de gelificação ionotrópica na presença de íons bivalentes, como Ca2+, formando hidrogéis e, por isso, tem sido utilizado para produzir sistemas de liberação

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Incorporação de fonte de nitrogênio em partículas de PVA e alginato de sódio e estudo da influência de ciclos de congelamento/descongelamento na caracterização do produto controlada de partículas de vários fármacos, proteínas e até mesmo células[1,6,9]. Muitos estudos vêm sendo feitos na combinação de polímeros, que constitui um método muito útil para a melhoria ou a modificação das propriedades físico-químicas dos materiais poliméricos. Este processo gera misturas estruturalmente diferentes, afetando as propriedades mecânicas, a morfologia, a permeabilidade e degradação dos polímeros iniciais[10]. Outro procedimento de melhora de propriedades mecânicas, morfológicas e térmicas é o tratamento criogênico[11]. Os criogéis são formados como resultado do congelamento de polímeros, armazenamento no estado congelado por um tempo definido, e seguido de descongelamento[12]. Após os polímeros passarem por ciclos de congelamento/descongelamento (CC/D), há a formação de uma rede interpenetrante devido às ligações cruzadas, formando um emaranhamento de interações de hidrogênio entre as cadeias, o que é proporcionado pela aproximação das mesmas durante o processo de congelamento[13]. Este processo pode resultar na eliminação de água da matriz polimérica e redução do tamanho dos poros do material final[1]. Um material novo a ser incorporado em matriz polimérica é a ureia. Molécula orgânica, nitrogenada, de ampla utilização na agricultura como fertilizante, dentre outras aplicações[14]. A ureia possui ponto de fusão de 133 °C e a sua decomposição térmica, muito complexa, inicia-se em cerca de 140 °C[15]. A ureia pode adquirir uma maior estabilidade térmica (cerca de 150-160 °C) quando intercalada em caulinita ([Si4]Al4O10(OH)8)[16], um argilomineral muito abundante no Brasil. Por ser um composto formado por camadas, a caulinita tem a capacidade de acomodar moléculas de ureia entre suas lamelas, atraves de ligações de hidrogênio, o que garante uma maior estabilidade térmica à ureia[16]. Em contrapartida, está comprovado que produtos de intercalação de ureia em caulinita são liberados rapidamente na presença de água[16]. Pensando nisso, propõe-se a incorporação de ureia e caulinita intercalada com ureia, em partículas de PVA+AS através da técnica de gelificação ionotrópica seguido de criogenia, com o intuito de estudar a incorporação de ureia e a influência de CC/D neste processo.

2. Experimental Inicialmente foi realizada a intercalação de ureia (marca Proquímios) em caulinita, proveniente do Rio Capim (Pará – Brasil) através do processo mecanoquímico, através do qual são misturadas ureia (U) e caulinita (C) na proporção 1:4, e moídas a seco com almofariz e pistilo de ágata por 7 horas[17].

2.1 Elaboração das partículas de PVA+AS Preparou-se uma solução de PVA (marca Vetec, hidrólise de 86,5-89,5%) em água quente (80 °C) e em seguida adicionou-se AS (marca Proquímios, viscosidade 520-580) na proporção de 3:1 (6% de PVA e 2% de AS)[1], a mistura ficou em agitação magnética por 12 horas. A solução de PVA+AS foi aquecida a aproximadamente 80 °C e gotejada através de uma bomba peristáltica da Tecnal, modelo Polímeros, 25(6), 606-613, 2015

TE‑BP-01-MINI, a 9,5 mL/min, em solução de CaCl2 a 5%[1] estando esta, em banho de gelo e sob agitação magnética[8]. As partículas formadas foram mantidas na solução de CaCl2 por 3 horas[18]. Em seguida, as partículas foram separadas da solução de CaCl2 e parte delas, colocadas em estufa a 60 °C por 48 horas, resultando nas partículas de PVA+AS S CC/D. Outra parte das partículas foi submetida ao congelamento a – 18 °C por 48 horas e descongeladas a 25 °C por 4 horas[1], o que corresponde a um CC/D. Após passarem por CC/D, as partículas também foram secas em estufa a 60 °C por 48 horas. Foram geradas partículas com 2 CC/D e com 4 CC/D.

2.2 Elaboração das partículas de PVA+AS+CI e PVA+AS+U As partículas de PVA+AS com caulinita intercalada (PVA+AS+CI) e de PVA+AS com ureia pura (PVA+AS+U) foram elaboradas através da mesma metodologia descrita acima. Para as partículas de PVA+AS+CI foi adicionado CI à solução polimérica na proporção de 3:1:0,48, o que corresponde a 2,14% de ureia em relação a massa seca de reagentes. Já para as partículas de PVA+AS+U foi adicionado ureia pura, na proporção de 3:1:0,18, o que corresponde a 4,31% de ureia em relação a massa seca de reagentes. Estas quantidades de CI e U foram definidas com base em testes realizados para verificar a consistência da solução formada, de maneira que concentrações mais elevadas de CI geram soluções muito espessas o que dificulta o gotejamento, pois obstruem as mangueiras. As soluções foram gotejadas a 8,0 mL/min para PVA+AS+CI, e a 9,5 mL/min para PVA+AS+U.

2.3 Análise elementar Foi realizada num analisador elementar Elemental Analyser 2400 CHN Perkin-Elmer Series II – USA, utilizando em torno de 1,0 mg da amostra pesado em balança analítica de 6 casas decimais. 2.4 Especroscopia no Infravermelho com Transformada de Fourier (FTIR) As análises de FTIR foram realizadas em espectrômetro Perkin Elmer, modelo Frontier, utilizando partilhas de KBr obtidas por prensagem juntamente com a amostra. Para cada espectro são somadas 8 varreduras com resolução de 2 cm–1, na região de 4000 a 400 cm–1.

2.5 Difratometria de Raios X (DRX) As amostras foram analisadas num difratômetro Bruker utilizando radiação Cu Kα (λ=1,5418 Å), com ângulos de varredura de 3 a 70° de 2θ, com passo de 0,05° de 2θ.

2.6 Análise Termogravimétrica (TG) e Termogravimetria Derivada (DTG) Para a Análise Térmica foi utilizado um analisador térmico TA Instruments, modelo SDT Q600. As amostras foram aquecidas a uma razão de 10 °C.min–1 de 30 a 800 °C, sob atmosfera estática de ar com fluxo de 50 mL.min–1. 607

Nardi, S. Q. W., Teixeira, S. D., & Parabocz, C. R. B.

3. Resultados e Discussão Na Figura 1 são apresentadas imagens das partículas de PVA+AS+U e PVA+AS+CI, ambas com 2 CC/D. Observa-se elevada semelhança entre as duas formulações de partículas, ambas com formato de “gota” devido a alta viscosidade das soluções que passaram pelo processo de gotejamento em solução de CaCl2.

3.1 Análise elementar As porcentagens finais de ureia nas partículas de PVA+AS+CI e PVA+AS+U apresentadas na Tabela 1. As partículas de PVA+AS também foram analisadas, mas foi confirmada a ausência de nitrogênio, como esperado. Observa-se que nem toda a ureia adicionada foi incorporada nas partículas. Essa perda nas partículas S CC/D, tanto nas de PVA+AS+U, quanto nas de PVA+AS+CI, pode ter ocorrido no processo de produção das partículas, pois as mesmas são feitas através do gotejamento da solução polimérica em solução aquosa de CaCl2. Como a ureia é altamente solúvel em água, no momento em que as partículas são formadas, parte da ureia pode ter sido lixiviada pela solução. Essa redução da quantidade de ureia incorporada também é observada com o aumento dos CC/D, fato que é devido à eliminação de água no processo de descongelamento das partículas. Quando a partícula é descongelada a água e/ou soluções aquosas presentes na estrutura, são eliminadas devido a aproximação e emaranhamento das cadeias de PVA[1]. Neste processo, a ureia por ser muito solúvel em água, pode ter sido carregada para fora da matriz polimérica,

Figura 1. Fotografia de partículas de (A) PVA+AS+U com 2 CC/D e (B) PVA+AS+CI com 2 CC/D. Tabela 1. Porcentagens finais de ureia nas partículas de PVA+AS+CI e PVA+AS+U S CC/D, com 2 CC/D e com 4 CC/D. Tipo de partícula PVA+AS+U

PVA+AS+CI

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S CC/D 2 CC/D 4 CC/D S CC/D 2 CC/D 4 CC/D

Porcentagem final de ureia na partícula (%) 0,986 0,761 0,407 0,504 0,311 0,204

sendo eliminada da partícula, gerando assim, uma diminuição da porcentagem de ureia incorporada.

3.2 Especroscopia no Infravermelho com Transformada de Fourier (FTIR) A Figura 2A apresenta os espectros de infravermelho de amostras de caulinita, ureia, caulinita intercalada com ureia, AS e PVA. Observam-se para a caulinita pura, bandas relativas à deformação axial das hidroxilas externas em 3698 cm–1, e das hidroxilas internas em 3620 cm–1[17]. Para a ureia, observam-se duas bandas, em 3445 e 3348 cm–1, provenientes da deformação axial simétrica e assimétrica de N-H, respectivamente[19]. Após a intercalação de ureia na caulinita, a principal alteração ocorre no deslocamento das duas bandas características da ureia, para 3505 e 3388 cm–1. Essas novas bandas são associadas à ligação da molécula de ureia às hidroxilas externas da lamela da caulinita, o que é uma evidência do processo de interação[17]. Tanto no espectro de PVA, quanto no de AS (Figura 2A), são observadas as bandas na região de 3400 cm–1, proveniente da deformação axial de OH associado, e em 2900 cm-1, devido a deformação axial da ligação C-H[1,2,19]. Para o espectro de AS, observa-se ainda, a deformação do íon –COO- que dá origem a duas bandas, uma em 1415 cm–1 e outra de maior intensidade em 1610 cm-1, provenientes das deformações axiais simétrica e assimétrica, respectivamente[1,2,8,19,20]. Para as partículas de PVA+AS (Figura 2B), observa-se que o emaranhamento das cadeias que ocorre no processo de elaboração das partículas não altera significativamente as estruturas dos polímeros envolvidos, pois os espectros das amostras de partículas apresentam as bandas nas mesmas regiões. As bandas nas regiões de 1415 e 1610 cm–1, provenientes da deformação simétrica e assimétrica do grupo funcional carboxílico, embora deslocadas para 1432 e 1630 cm–1, ainda permanecem. Este deslocamento para frequências maiores é devido a reticulação com íons Ca2+, os quais interagem com os grupos funcionais carboxílicos unindo as cadeias de alginato, este resultado também foi encontrado na literatura[1]. Através dos espectros de partículas de PVA+AS+CI (Figura 2C), pode-se observar que as mesmas conservaram as mesmas bandas características dos polímeros utilizados. Apresentam também, uma das bandas vistas anteriormente no espectro de infravermelho da CI (Figura 2Ab), na região de 3696 cm–1, relativa à deformação axial das hidroxilas externas. A banda em 826-827 cm–1 observada nas partículas de PVA+AS e PVA+AS+U, não é observada nas partículas de PVA+AS+CI, pois fica sobreposta pela mesma região de picos intensos da caulinita. Muitos picos característicos da CI ficaram encobertos pelos picos característicos dos polímeros, o que indica que a caracterização por FTIR não é suficiente para a comprovação da incorporação de nitrogênio nas partículas. Nas partículas de PVA+AS+U (Figura 2D), a banda característica da deformação axial de O-H (3400 cm–1) é deslocada para frequências mais baixas com o aumento dos CC/D. Este deslocamento é decorrente do aumento de interações de hidrogênio intermoleculares, as quais interferem na força de ligação covalente-polar O-H, enfraquecendo-a, o que consequentemente interfere no sinal de FTIR detectado Polímeros , 25(6), 606-613, 2015

Incorporação de fonte de nitrogênio em partículas de PVA e alginato de sódio e estudo da influência de ciclos de congelamento/descongelamento na caracterização do produto

Figura 2. Espectros de infravermelho de amostras de (Aa) caulinita pura, (Ab) caulinita intercalada com ureia, (Ac) ureia pura, (Ad) alginato de sódio e (Ae) PVA; partículas de PVA+AS com (Ba) 4 CC/D, (Bb) 2 CC/D e (Bc) S CC/D; partículas de PVA+AS+CI com (Ca) 4 CC/D, (Cb) 2 CC/D e (Cc) S CC/D; e partículas de PVA+AS+U, com (Da) 4 CC/D, (Db) 2 CC/D e (Dc) S CC/D.

de forma a deslocar a banda para frequências menores[20]. Esse fato ocorre apenas nas partículas com ureia, a qual possui grandes possibilidades de se acomodar entre as cadeias poliméricas, devido ao seu tamanho reduzido, e formar ligações de hidrogênio com as hidroxilas dos polímeros, o que pode ter provocado o deslocamento da banda da região de 3400 cm–1. Este deslocamento se intensifica com o aumento dos CC/D, através dos quais há uma aproximação ainda maior das cadeias poliméricas, formando uma rede interpenetrante devido às interações de hidrogênio e consequente emaranhamento das cadeias de PVA[1,13]. Este aumento de reticulação aumenta a aproximação das cadeias poliméricas com as moléculas de ureia, facilitando a formação de interações de hidrogênio. Outra característica importante é a diminuição da intensidade da banda na região de 1630 cm–1 com o aumento dos CC/D e o surgimento de duas novas bandas em 1667 e 1597 cm–1, nas partículas com 4 CC/D. A banda na Polímeros, 25(6), 606-613, 2015

região de 1630 cm-1 é proveniente da deformação assimétrica de –COO–, o qual pode estar sofrendo influencia da presença de ureia na estrutura com o aumento da reticulação pelos CC/D. O surgimento das duas novas bandas pode estar relacionado com a formação de amidas secundárias, as quais apresentam sinal de dobramento N-H na região de 1640-1550 cm–1[20]. Outra possibilidade é a formação de oximas (R-CH=N-O-H) as quais possuem absorção de C=N na região de 1690-1640 cm–1, e ainda uma absorção larga de O-H em 3650-2600 cm–1[20]. Este último fator, enquadra a absorção detectada em 3623 e 3644 cm–1 nas partículas de PVA+AS+U com 2 e 4 CC/D, respectivamente[20]. As partículas de PVA+AS+U foram as que apresentaram maior influencia dos CC/D na sua reticulação e na sua estrutura, devido a contribuição da ureia na formação de interações de hidrogênio, possivelmente em vários pontos das cadeias poliméricas. 609

Nardi, S. Q. W., Teixeira, S. D., & Parabocz, C. R. B. 3.3 Difratometria de Raios X (DRX) Através da Figura 3A observa-se o pico de difração da caulinita em aproximadamente 12° de 2θ, que é característico da reflexão 001[16,21]. Já a ureia apresenta difração em aproximadamente 22° de 2θ. Após a intercalação, observa-se

o surgimento de uma difração em aproximadamente 8° de 2θ, característico da caulinita intercalada. No difratograma de raios X de AS (Figura 3Ba) observam‑se três picos característicos em, aproximadamente, 13,6° de 2θ[1,4], 21,2° de 2θ[4,7] e 37,8° de 2θ[7] observados

Figura 3. Difratogramas de Raios X de amostras de (Aa) ureia, (Ab) caulinta e (Ac) caulinita intercalada com ureia; (Ba) alginato de sódio e (Bb) PVA; partículas de PVA+AS com (Ca) 4 CC/D, (Cb) 2 CC/D e (Cc) S CC/D; partículas de PVA+AS+CI com (Da) 4 CC/D, (Db) 2 CC/D e (Dc) S CC/D; e partículas de PVA+AS+U com (Ea) 4 CC/D, (Eb) 2 CC/D e (Ec) S CC/D. 610

Polímeros , 25(6), 606-613, 2015

Incorporação de fonte de nitrogênio em partículas de PVA e alginato de sódio e estudo da influência de ciclos de congelamento/descongelamento na caracterização do produto também, em outros trabalhos. Para o PVA (Figura 3Bb) observa-se picos de difração em 11,5° de 2θ, 19,5° de 2θ, 22,4° de 2θ e 40,5° de 2θ, valores semelhantes aos encontrados em outros estudos[4,7,22]. Os difratogramas de amostras de PVA+AS com 4 e com 2 CC/D (Figura 3Ca,b) apresentam características muito semelhantes. Os picos de difração característicos do PVA apresentam-se mais largos devido à redução da cristalinidade. Isso se deve a possível formação de interações de hidrogênio entre os grupos hidroxila do PVA e grupos hidroxila ou carboxila do AS[2,4]. Ocorre a redução da cristalinidade com o aumento dos CC/D, devido ao aumento das ligações cruzadas[23]. O não aparecimento dos picos na região de 32° de 2θ e 46° de 2θ nas amostras com CC/D pode ser devido à mudança de disposição das cadeias, ocorrida durante este processo. Com relação às partículas de PVA+AS+CI (Figura 3D), pode-se observar que para às partículas sem CC/D existe um pico na região de 8,5° de 2θ, região característica da presença de caulinita intercalada. Este pico existe apenas para as partículas sem CC/D, e para as partículas com CC/D surge novamente um pico na região de 11° de 2θ a 12° de 2θ, região característica da presença de caulinita pura. Isso indica que após os CC/D, parte da ureia foi removida dos espaçamentos interlamelares da caulinita restando quantidades indetectáveis por DRX. Este resultado é comprovado pela Análise Elementar, que comprova uma diminuição da quantidade de nitrogênio com o aumento dos CC/D. Para as partículas de PVA+AS+U (Figura 3E) não é detectada a presença de ureia na estrutura, a qual apresentaria um pico na região de 22° de 2θ, apesar da Análise Elementar comprovar a existência de ureia nestas partículas. Assim como em FTIR, a técnica de difratometria de Raios X mostrou características diferentes para as partículas de PVA+AS+U. Para as partículas com CC/D surgem os picos nas regiões de 32° de 2θ e 46° de 2θ, os quais podem estar relacionados ao fato a ureia estar ligada as cadeias poliméricas através de interações de hidrogênio, o que já foi sugerido através dos resultados de FTIR, pois houve um deslocamento da banda característica da deformação axial de O-H (Figura 2D). Ainda pode estar relacionado com o surgimento dos dois novos picos no espectro de FTIR em 1667 e 1597 cm–1, para as partículas de PVA+AS+U com 4 CC/D (Figura 2Da), os quais podem ter sido gerados pela formação de amidas secundárias, ou oximas, como discutido anteriormente.

3.4 Análise Termogravimétrica (TG) e Termogravimetria Derivada (DTG) Todas as amostras de partículas, inclusive de amostras de C, U, CI e dos polímeros puros, foram submetidas à análise térmica e os resultados são apresentados na Tabela 2. Através dos resultados pode-se observar a temperatura inicial de decomposição de cada amostra, bem como a temperatura de decomposição máxima de cada etapa de decomposição. Para a CI, observa-se claramente que as três primeiras etapas de decomposição são referentes a ureia e a última etapa é referente a caulinita, indicando que houve a eliminação da matéria orgânica (ureia) sem a destruição da matriz em questão (caulinita)[17]. Resultados compatíveis a este trabalho foram Polímeros, 25(6), 606-613, 2015

encontrados na literatura para caulinita[17,21,24], ureia[15,17], caulinita intercalada[17,24], PVA[2,4,5] e AS[5,7,8]. A decomposição térmica das partículas ocorre em três etapas definidas pela DTG. As partículas de PVA+AS com CC/D apresentam uma temperatura inicial de decomposição maior do que as partículas S CC/D (188 °C > 195 °C > 196 °C). Este fato indica que os ciclos de CC/D contribuem para uma leve melhora da estabilidade térmica do produto. Já as temperaturas de pico da curva DTG não apresentam aumento regular com o aumento dos CC/D. Para as partículas de PVA+AS+CI observa-se um aumento na estabilidade térmica das partículas com o aumento dos CC/D, pois a temperatura de pico da curva de DTG da primeira etapa de decomposição aumenta com o aumento dos CC/D (226,3 °C > 235,8 °C > 240,0 °C). Outro indício de aumento de estabilidade térmica é que quanto maior o número de CC/D, maior é a temperatura final de decomposição. Este aumento da resistência térmica com o aumento do número de CC/D, nos diferente tipos de partículas, pode ser resultado da formação de um emaranhado de cadeias de PVA, unidas por interações de hidrogênio, que foram facilitadas pela aproximação das mesmas durante o congelamento[1]. Este fator dificulta a mobilidade das cadeias dos polímeros e a quebra de ligações. O aumento de estabilidade térmica com o aumento de CC/D também é observado para as partículas de PVA+AS+U. Isto é visível observando as temperaturas iniciais da primeira e terceira etapa de decomposição das partículas, as quais vão aumentando com o aumento dos CC/D. Este aumento também é observado nas temperaturas de decomposição da primeira etapa de decomposição. Outro indício é a temperatura final de decomposição, a qual aumenta com o aumento dos CC/D (714 °C > 735 °C > 742 °C). Em comparação com as partículas de PVA+AS+CI, as partículas de PVA+AS+U apresentam temperaturas de decomposição máxima maiores na primeira etapa de decomposição térmica. As temperaturas finais de decomposição são consideravelmente maiores para as partículas de PVA+AS+U, indicando maior lentidão na decomposição térmica para estas partículas. Este aumento de estabilidade com a adição de ureia pode estar relacionado ao fato de as moléculas de ureia estabelecerem interações de hidrogênio facilmente com as hidroxilas de grupamentos –COOH dos polímeros formadores das partículas. Já a caulinita, por apresentar estrutura volumosa, é acomodada na matriz polimérica com maior dificuldade, mesmo também apresentando grupos passíveis de interações de hidrogênio. O fato de a caulinita estar intercalada contribui para o aumento do espaço ocupado pela molécula na matriz polimérica, pois a intercalação aumenta o espaçamento basal entre as lamelas deste argilomineral. Portanto, pode-se dizer que dentre todas as partículas analisadas, as que apresentaram melhor estabilidade térmica foram as de PVA+AS+U, fato que é intensificado pela presença dos CC/D. Observa-se na Tabela 2 que praticamente todas as amostras apresentam massa residual em temperatura de 800 °C, e com o aumento dos ciclos de congelamento/descongelamento a massa residual diminui. Em contrapartida, a estabilidade térmica em cada etapa de degradação aumenta quando o número de CC/D é maior. 611

Nardi, S. Q. W., Teixeira, S. D., & Parabocz, C. R. B. Tabela 2. Valores dos intervalos de temperatura das etapas de decomposição (TG), temperatura de decomposição máxima de cada etapa de decomposição (DTG), perda de massa não cumulativa e total de amostras de C, U, CI, PVA, AS e das partículas de PVA+AS, PVA+AS+CI e PVA+AS+U, S CC/D, com 2 CC/D e com 4 CC/D.

380,0-800,0

Temperatura de decomposição máxima (pico em DTG) (°C) 494,4

130,0-205,0 205,0-250,0 250,0-300,0 300,0-370,0

175,5 244,7 295,4 365,0

75,87 10,25 10,87 1,643

CI (U) CI (U) CI (U) CI (C)

130,0-180,0 180,0-235,0 235,0-250,0 350,0-800,0

161,0 212,3 245,0 494,4

6,554 11,10 1,070 13,43

65,08

PVA

214,0-379,0 379,0-455,0 455,0-533,0

326,8 415,0 479,6

45,13 32,29 20,02

0,0

200,0-285,0 535,0-578,0

241,7 576,5

36,32 13,96

20,97

PVA+AS S CC/D

188,0-371,0 371,0-486,0 486,0-800,0

223,6 432,6 518,4

28,85 14,58 16,47

29,63

PVA+AS 2 CC/D

195,0-379,0 379,0-477,0 477,0-800,0

236,2 434,9 498,5

33,63 16,89 21,96

23,16

PVA+AS 4 CC/D

196,0-376,0 376,0-478,0 478,0-800,0

226,1 431,1 519,1

29,07 13,48 20,60

17,97

PVA+AS+CI S CC/D

183,0-366,0 379,0-477,0 477,0-623,0

226,3 432,8 505,6

26,06 10,99 14,08

33,61

PVA+AS+CI 2 CC/D

183,0-362,0 362,0-470,0 470,0-636,0

235,8 428,7 494,8

26,30 11,72 15,18

31,49

PVA+AS+CI 4 CC/D

191,0-367,0 367,0-483,0 483,0-642,0

240,0 433,5 520,7

32,02 17,74 17,36

23,86

PVA+AS+U S CC/D

195,0-366,0 366,0-473,0 473,0-714,0

243,7 428,7 506,1

26,85 10,54 16,46

29,95

PVA+AS+U 2 CC/D

200,0-370,0 370,0-477,0 477,0-735,0

245,6 430,3 510,2

31,35 13,91 18,08

22,82

PVA+AS+U 4 CC/D

200,0-367,0 367,0-471,0 471,0-742,0

251,2 429,8 506,1

29,10 13,01 17,78

26,61

Amostra

Intervalo aproximado de temperatura (TG) (°C)

4. Conclusões A incorporação de ureia em matriz de PVA+AS foi possível, sendo comprovado por análise elementar e DRX que a eficiência de incorporação diminui com o aumento do número de CC/D. As partículas de PVA+AS+U foram 612

Perda de massa não cumulativa (%)

Perda de massa total (%)

14,87

83,28 0,527

as que obtiveram maior alterações na matriz polimérica, maior estabilidade térmica, e elevaram a decomposição da ureia imobilizada a 200 °C, tendo em vista que a ureia pura se decompõe em aproximadamente 130 °C. Além disso, apresentam a vantagem de serem mais viáveis economicamente, pois não apresentam caulinita em sua formulação. Polímeros , 25(6), 606-613, 2015

Incorporação de fonte de nitrogênio em partículas de PVA e alginato de sódio e estudo da influência de ciclos de congelamento/descongelamento na caracterização do produto

5. Agradecimentos A CAPES pela concessão da bolsa de mestrado, a UTFPR e ao professor Fauze Jaco Anaissi da UNICENTRO, Campus Guarapuava – PR, pelas análises de DRX.

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