Page 1

DRIVING INNOVATION FORWARD Polímeros

New materials like plastics are advancing vehicle technologies. But getting the best results is a big challenge. To assist, SABIC offers industry-leading expertise in designing with a wide range of engineered materials. Across the entire vehicle. Because no matter what obstacles may hold our customers back, we’re there with ‘Chemistry that MattersTM’ to help them drive forward.

SABIC.com

VOLUME XXVI - N° 1 - JAN/FEV/MAR - 2016

© 2016 Copyright SABIC. All rights reserved.


http://dx.doi.org/10.1590/0104-1428.2601

Editorial

Dear Author, Presently, the resources that the journal Polímeros: Ciência e Tecnologia receives from the funding agencies do not cover all the expenses to maintain the quality of its content and the periodicity. Additionally, Polímeros is as a fully open access journal. These expenses include printing, digital maintenance, personal, postal mail, desktop editing, etc. ABPol covers part of these expenses; however, the publication costs have increased considerably in the last years. To cope with this, we are campaigning to attract donors. Therefore, we are kindly asking the authors of accepted papers to contribute with US $ 25.00 (nowadays R$ 100,00) per published page (like many international journals around the world propose today). We ask the authors to help with any possible amount, although it is also convenient to emphasize that the non‑payment of this contribution will not prevent the article from publication. All articles will go through the same normal editorial procedures. Payment does not guarantee publication priority. This campaign is only an alternative to maintain our periodicity. Please think carefully to enhance your involvement with Polímeros. If you agree, see next page for instructions to carry on the fee payment. i) For Authors or Institutions in Brazil we will provide a receipt, described as “copydesk payment - taxa de editoração”. ii) For overseas, a special receipt will be provided after confirming the paypal payment.

We thank very much for your comprehension. Sincerely,

The Editorial Council Polímeros: Ciência e Tecnologia

Polímeros, 26(1), 2016    

E1

E E E E E E E E E


P o l í m e r o s - N º 1 - V o l u m e X X V I - J a n / F ev / M a r - 2 0 1 6 - ISS N 0 1 0 4 - 1 4 2 8 - ISS N 1 6 7 8 - 5 1 6 9

E

( ve r s ã o

eletrônica)

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

X

Polímeros P r e s i d en t e

do

Conselho Editorial

Marco-Aurelio De Paoli (UNICAMP/IQ)

P

Membros

do

Conselho Editorial

E

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

Comitê Editorial Sebastião V. Canevarolo Jr. – Editor

Membros

do

Comitê Editorial

Adhemar C. Ruvolo Filho Alain Dufresne Bluma G. Soares César Liberato Petzhold João B. P. Soares José António C. Gomes Covas José Carlos C. S. Pinto Regina Célia R. Nunes Richard G. Weiss Rodrigo Lambert Oréfice

D

Produção

e

Assessoria Editorial

www.editoracubo.com.br

I

“Polímeros” é uma publicação da Associação Brasileira de Polímeros Rua São Paulo, nº 994 13560-340 - São Carlos, SP, Brasil Fone/Fax: (16) 3374-3949

E

e-mails: abpol@abpol.org.br / revista@abpol.org.br http://www.abpol.org.br Data de publicação: Março de 2016

Apoio:

N T

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

E

Trimestral v. 26, nº 1 (Jan./Fev./Mar. 2016) ISSN 0104-1428 ISSN 1678-5169 (versão eletrônica)

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

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

Polímeros, 26(1), 2016


E

Polímeros Seção Editorial

X

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

S e ç ã o T é cn i c a Synthesis and applications of polystyrene-block-poly(N‑vinyl-2-pyrrolidone) copolymers Marcelo Alexandre de Farias and Maria do Carmo Gonçalves.......................................................................................................................... 1 Vaishnav Kiran and Bharti Gaur....................................................................................................................................................................... 11

P

Curing and thermal behavior of epoxy resins of hexafluoro - bisphenol –A and bisphenol-A Cardanol-based thermoset plastic reinforced by sponge gourd fibers (Luffa cylindrica) André Leandro da Silva, Lucas Renan Rocha da Silva, Isabelle de Andrade Camargo, Deuber Lincon da Silva Agostini, Derval dos Santos Rosa, Diego Lomonaco Vasconcelos de Oliveira, Pierre Basílio Almeida Fechine and Selma Elaine Mazzetto............... 21

Effect of the hardener to epoxy monomer ratio on the water absorption behavior of the DGEBA/TETA epoxy system PF/CLAY hybrid materials: a simple method to modulate the optical properties Marcio Chao Chen Em, Camila Gouveia Barbosa, Laura Oliveira Péres and Roselena Faez......................................................................... 38

E

Ayrton Alef Castanheira Pereira and José Roberto Moraes d’Almeida............................................................................................................ 30

Thermal and catalytic pyrolysis of plastic waste Débora Almeida and Maria de Fátima Marques............................................................................................................................................... 44

Mechanical and thermomechanical properties of polyamide 6/Brazilian organoclay nanocomposites

Correlation between stabilizer consumption and degree of polymerization of thermally upgraded paper aged in insulating natural ester and insulating mineral oil

D

Renê Anisio da Paz, Amanda Melissa Damião Leite, Edcleide Maria Araújo, Vanessa da Nóbrega Medeiros, Tomás Jeferson Alves de Melo and Luiz Antônio Pessan.................................................................................................................................................................................... 52

Larissa Mildemberger, Mario Carlos Andreoli, Guilherme Cunha da Silva, Heloisa Nunes da Motta, Joseane Valente Gulmine and Marilda Munaro............................................................................................................................................................................................................... 61

Poly(lactic acid)/thermoplastic starch sheets: effect of adipate esters on the morphological, mechanical and barrier properties

TLC/IR (UATR) off-line coupling for the characterization of additives in EPDM rubber compositions

I

Marianne Ayumi Shirai, Juliana Bonametti Olivato, Ivo Mottin Demiate, Carmen Maria Olivera Müller, Maria Victória Eiras Grossmann and Fabio Yamashita.......................................................................................................................................................................................... 66 Denis Damazio, Eunice Aparecida Campos, Milton Faria Diniz, Elizabeth da Costa Mattos and Rita de Cássia Lazzarini Dutra............... 74

Otimização do processo de dispersão de nanotubos de carbono em poliuretano termorrígido Magnovaldo Carvalho Lopes, João Paulo Campos Trigueiro, Vinicius Gomide de Castro, Rodrigo Lassarote Lavall and Glaura Goulart Silva.................................................................................................................................................................................................................... 81

E

Investigação do efeito do tempo de exposição à temperatura ambiente e ao tempo de estocagem de um filme adesivo estrutural de resina epoxídica Ana Carolina Teixeira Neves da Silva, Fernanda Guilherme, Vanesa Mitchell Ferrari e Paulo Eduardo Ferrari........................................... 92

N

Capa: Imagens de MO da superfície dos compósitos contendo 0,5% em massa de: (a) MWCNTs não modificados e (b) MWCNTs modificados. Imagens de MEV da superfície da fratura dos nanocompósitos contendo 0,5% em massa de: (c) MWCNT não modificados e (d) MWCNTs modificados e imagens de MET de cortes de criomicrotomia dos nanocompósitos contendo 0,5% em massa de: (e) MWCNTs não modificados e (f) MWCNTs modificados. Elaboração artística Editora Cubo.

T

E

Polímeros, 26(1), 2016

E3


I N F O R M E S

A polymer sponge to repair broken backs

E N O T Í C I A S

Researchers at the Mayo Clinic have developed a novel spinal graft that automatically “grows” to the requisite size and shape when implanted. The spongy polymer isn’t meant to be a formal replacement like the 3D printed neck bones recently installed by a team from the Prince of Wales Hospital in Sydney. Instead, it’s designed to act as a bone graft -- a biodegradable scaffold through which a cancer patient’s own bones can regrow after surgery. See, for whatever biological reason, when cancer metastasizes -- that is, when it spreads through the body from the original tumor -- it predominantly tends to settle in the spinal column. Doctors can cut out the infected bone tissue (or flat-out replace it as they did in the Sydney case) but that leaves large gaps in the spine. Normally, doctors would either have to open the chest cavity and access the spine from far side (which entails a lengthy recovery and high probability of complications) or they’d make a small incision in the neck/back and inject expandable titanium rods into the bone gap (which is super expensive because titanium). This new technique combines the easy access and short recovery of the titanium rod method with the low cost of the open chest operation. Doctors simply cut a small hole in the patient’s neck/back and inject a hydrogel polymer into the bone gap much the same way they would a titanium rod. This polymer absorbs fluids from within the wound and grows to fill the gap. Doctors control how far the polymer expands in any specific direction by first inserting a “cage” -- basically a pre-expanded shell that the polymer fills in as it spreads. Think of it as the wooden frame that keeps a freshly-poured concrete sidewalk in place until it hardens. Once the polymer fills in the cage, which takes 5 to 10 minutes on average, it will set and harden into a viable prosthetic. From there, surrounding bone tissue grows into and through the polymer, reinforcing and cementing it in place. The Mayo team plans to begin initial cadaver-based tests of the new method in the coming months. Should those succeed, trials on living people will follow in the next couple of years. Source: American Chemical Society

E4

Degradable polymer film a new tool in modern farming Plastic film the thickness of a human hair has the potential to give some farmers a break from the perils of the harsh and unpredictable Australian climate. The film, made from degradable polymer, was developed by a team in partnership with the Co‑operative Research Centre (CRC) for Polymers and is laid over crops during the seed-planting process. CRCP CEO Ian Dagley said the film uses heat, light and ground moisture to stimulate plant growth of both native trees and crops. “It is rolled out at the time the farmer is planting the seed and mechanically buried at the edges, creating a temporary greenhouse,” Mr Dagley said. “The film is placed in such a way that the seed is able to germinate and grow, where the degradation of the film is timed to be weak at the time the plant has grown to the point where it is trying to break out – typically about 10cm. “From the time you plant the seed until the time the plant breaks through... the soil moisture that is there and transpired moisture from the plant is recycled in that closed system.” It takes around a month for the polymer film to degrade to the point of being weak enough for plants to break through, with ongoing testing of the material for different crop applications. “We have timed it so that the material below the ground also degrades, meaning the farmer can plow the field next season and the film is not a problem,” Mr Dagley said. Trials of the product have shown more uniform germination and earlier growth, with potentially significant benefits for Queensland farmers living with severe water shortages. “You’ve trapped that moisture at the start of the season, so if you are going into a season that was very dry, the fact the crop was further developed early on might provide you with a benefit and potentially a better yield,” Mr Dagely said. “In the case of vegetable crops, hypothetically it may be a case of using the film to plant earlier and getting crops to market earlier.” The research has shown no impact on soil quality or food safety, and the product is being adapted to Australian conditions and farming practices in a series of field trials run across four states. The Bureau of Meteorology has said areas of serious to severe rainfall deficiencies to remain largely unchanged to the end of February in Queensland’s central interior, while drought conditions in the north have become more severe and slightly more widespread. Source: Brisbane Times

Polímeros, 26(1), 2016


June

September

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

XIIth French-Romanian Polymer Meeting Date: 5-7 September 2016 Local: Sophia Antipolis – France Website: https://www.sciencesconf.org/browse/ conference/?confid=2932

COMPLAST Kenya Plast 2016 Date: 8–10 June 2016 Local: Nairobi - Kenia Website: http://www.kenyaplast.in/ 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 POLYMAT 2016 - Silesian Mettings on Polymer Materials Date: 27–28 June 2016 Local: Zabrze - Poland Website: https://cmpw-pan.edu.pl/polymat2016/

July COMPLAST Plastics Myanmar 2016 Date: 08–10 July 2016 Local: Yangon - Myanmar Website: http://www.plastics-myanmar.in/ 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 3rd Brazilian Conference on Composite Materials (BCCM-3) Date: 28-31 August 2016 Local: Gramado - RS Website: http://www.bccm.com.br

Polycondensation 2016 Date: 11–15 September 2016 Local: Moscow / St Petersburg - Russian Website: http://www.polycondensation2016.ac.ru/index.php/en/ Polyolefin Additives – 2016 Date: 13–15 September 2016 Local: Vienna - Austria Website: http://www.amiplastics.com/events/event?Code=C743 PLASTEC Minneapolis Date: 21–22 September 2016 Local: Minnesota - USA Website: http://plastecminn.plasticstoday.com/ Organic Semiconductors Date: 22–25 September 2016 Local: Dubrovnik - Croatia Website: http://www.zingconferences.com/conferences/organicsemiconductors/ 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 Polymer Foam – 2016 Date: 8–10 November 2016 Local: Cologne - Germany Website http://www.amiplastics.com/events/event?Code=C752 Expoplast 2016 Date: November 30 - December 1, 2016 Local: Québec - Canada Website: http://expoplast.plasticstoday.com/

December Fire Resistance in Plastics 2016 Date: 6-8 December 2016 Local: Cologne - Germany Website: http://www.amiplastics-na.com/events/Event. aspx?code=C719&sec=7121

Polímeros, 26(1), 2016 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

E6

Polímeros, 26(1), 2016


Associados da ABPol As nossas boas vindas...

Ao novo Sócio Patrocinador SGS Labmat. Agradecemos o valioso apoio! Luiz Antonio Pessan Presidente

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. Formax Quimiplan Componentes para Calçados Ltda. Imp. e Export. de Medidores Polimate Ltda. Innova S/A Instituto de Aeronáutica e Espaço/AQI Jaguar Ind. e Com. de Plásticos Ltda 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, 26(1), 2016 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/10.1590/0104-1428.2066

Synthesis and applications of polystyrene-blockpoly(N‑vinyl-2-pyrrolidone) copolymers Marcelo Alexandre de Farias1* and Maria do Carmo Gonçalves1 Institute of Chemistry, Universidade Estadual de Campinas – UNICAMP, São Paulo, SP, Brazil

1

*marcelo.adf@gmail.com

Abstract This work describes the synthesis and applications of amphiphilic polystyrene-block-poly(N-vinyl-2-pyrrolidone) (PS-b-PVP) copolymers as a silver and silica nanoparticle surface modification agent. The synthesis of PS-b-PVP was carried out using controlled/living radical polymerization techniques. The synthesis of the block copolymers was confirmed by gel permeation chromatography and hydrogen nuclear magnetic resonance, presenting a polydispersity index of around 1.4 and number average molecular weight ranging between 10,000-14,000 g mol-1. The PS-b-PVP copolymers were applied as a silver nanoparticle (AgNP) stabilizing agent. These nanoparticles were produced by a single step and presented an 11 ± 1 nm diameter. Furthermore, the PS-b-PVP copolymers were also applied as a silica nanoparticle (SiO2NP) surface modification agent. The SiO2NP were synthesized by the Stöber method presenting a 72 ± 9 nm diameter. The SiO2NP surface modification by adsorption of PS-b-PVP caused the formation of a 5 ± 1 nm thick polymeric layer, providing the SiO2NP with a hydrophobic surface character. The structural and chemical characteristics shown by PS-b-PVP copolymers highlights their versatility for several applications, such as: water-in-oil emulsifier, stabilizing or coupling agents between inorganic particles and polymeric matrices. Keywords: amphiphilic block copolymers, controlled radical polymerization, silica nanoparticles, silver nanoparticles, surface modification agent.

1. Introduction Currently, using controlled/living radical polymerization (CRP) techniques, efforts in new polymer structures are focused on producing amphiphilic block copolymers[1-3]. These macromolecules are made up of two chemically different homopolymer blocks (A and B) combined as AB or ABA: one being hydrophilic and the other being hydrophobic. Polystyrene is a classic example of a hydrophobic glassy polymer frequently synthesized via CRP techniques[4–7]. On the other hand, an example of hydrophilic polymer is poly(N-vinyl-2-pyrrolidone) (PVP) which has been synthesized via CRP in the last few years[8-10]. PVP is an important building block and has attracted significant attention from the biomedical field because it is biocompatible and non-toxic, being an excellent candidate to replace PEO in certain biomaterial applications[8]. Furthermore, PVP is also known for its ability to easily form complexes with metals[11]. Amphiphilic block copolymers, when dissolved in a selective solvent, tend to self-assemble producing core-shell micelles. These micelles are capable of encapsulating metallic particles obtained from their metal salts[12]. In this case, the core block is able to entrap particles by complexation or association, and the shell block provides a hydrophobic or hydrophilic character for the colloidal nanoparticle. Furthermore, several experimental and theoretical studies have been reported in relation to the surface modification by block copolymers in order to promote specific characteristics[12–15]. Zhang et al.[16] used poly(ethylene oxide)-block-poly(methyl methacrylate) (PEO-b-PMMA) copolymer as a template for the synthesis of silver nanowires in an aqueous solution. In this system, the PMMA block reduces silver ions and PEO block promotes the nanoparticle dispersion in water. Li et al.[17]

Polímeros, 26(1), 1-10, 2016

modified the surface of Fe3O4 magnetic nanoparticles with amphiphilic poly(tert-butyl methacrylate)-block-poly(glycidyl methacrylate) (PtBMA-b-PGMA) in order to improve the microwave-assisted extraction of polycyclic aromatic hydrocarbons (PAH) in environmental water. The PtBMA block provides a highly hydrophobic property for the Fe3O4 magnetic nanoparticles responsible for adsorbing PAH. On the other hand, the PGMA block is responsible for the immobilization of the block copolymer on the magnetic nanoparticles. The authors observed that the extraction of PAH is improved when Fe3O4 magnetic nanoparticles are encapsulated by PtBMA-b-PGMA. Amphiphilic copolymers of polystyrene-block-poly(N‑vinyl2-pyrrolidone) (PS-b-PVP) have been reported employing the following radical polymerizations: pseudo-living[18], atom transfer radical polymerization (ATRP)[9,19], ATRP followed by reversible addition−fragmentation chain-transfer polymerization (RAFT)[20], RAFT[21], RAFT followed by ATRP[22], nitroxide-mediated polymerization (NMP)[23,24], organostibine-mediated[25] and organogermanium-mediated polymerizations[26]. Although PS-b-PVP copolymers have been synthesized by several methods, some did not provide suitable PS-b-PVP amphiphilic copolymers due to the presence of homopolymers as impurities and high Ɖ values (i.e. > 2.0)[18,26] and others involve complex experimental steps in order to produce intermediate species (e.g. synthesis of mediators and chain transfer agents)[21,25]. Furthermore, the PS-b-PVP copolymer is not commercially available and has potential applications in polymer science (e.g. stabilization, dispersibility, emulsifier, coupling agents), which are not sufficiently exploited.

1

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


de Farias, M. A., & Gonçalves, M. C. In this work, we describe the synthesis of amphiphilic PS-b-PVP copolymers by RAFT and their application as a silver and silica nanoparticle stabilizing or a surface modification agent, respectively. Although RAFT technique has the advantage of being applicable to a much wider variety of monomers than the ATRP technique, usually the synthesis of RAFT end-functionalized initiators (chain transfer agents) is not so simple[21]. For this reason, styrene was initially synthesized in bulk by ATRP leading to polystyrene with bromine chain end functionality. Then, a suitable macro chain transfer agent (macro-CTA) was formed by performing a simple method[20,27] to convert the polystyrene ATRP product into the RAFT agent. N-vinyl-2pyrrolidone (VP) was subsequently used as a chain extensor of this RAFT agent to synthesize well-defined amphiphilic PS-b-PVP copolymers. Applying this methodology, the use of toxic solvents or complex experimental steps was avoided. The block copolymers produced were structurally characterized by hydrogen nuclear magnetic resonance (1H NMR) and gel permeation chromatography (GPC). Thermal properties were evaluated by differential scanning calorimetry (DSC). Furthermore, the PS-b-PVP copolymers were applied not only to stabilize in-situ prepared silver nanoparticles (AgNP) but were also used to modify the silica nanoparticle (SiO2NP) surface in order to show specific applications for these copolymers.

2. Materials and Methods 2.1 Materials Styrene (>99%), N-vinyl-2-pyrrolidone (VP, >99%), copper(I) bromide (CuBr, 98%), 5,5′-dimethyl-2,2′-dipyridyl (98%), potassium ethyl xanthogenate (96%), (1-bromoethyl) benzene (97%), dimethylacetamide (DMAc), tetraethyl orthosilicate (TEOS, 98%) and 2,2′-azobis(2-methylpropionitrile) solution (AIBN, 0.2 mol L-1 in toluene) were purchased from Aldrich. Acetone, ethanol, methanol, ammonium hydroxide solution (27%), silver nitrate (AgNO3), sodium hydroxide (NaOH), dimethylacetamide (DMAc) and tetrahydrofuran (THF) were purchased from Vetec-Sigma (Brazil).

2.2 Removal of inhibitors present in the monomers Equal volumes of styrene and aqueous NaOH solution (10% w/v) were vigorously stirred in a 500 mL separation funnel. Then, phase separation was observed and the bottom phase was discarded. The upper phase was washed three times with distilled water, always discarding the bottom phase. After, the styrene fraction was distilled at 40 ± 2 °C and the pure styrene was collected in a flask immersed in liquid nitrogen. The VP monomer was distilled under vacuum using a Vigreux column. The initial distillation temperature was 90 °C being gradually increase up to 100 °C.

2.3 Preparation of PS-Br homopolymer via ATRP The [monomer]:[initiator]:[CuBr]:[ligand] ratios used in this section, according to Hayes and Rannard[28], were 96:1:1:2.5. Additionally, an identical synthesis was carried out in the same ratio, except for the monomer to initiator ratio, which was 192:1. 2

Styrene (10.0 mL, 87.5 mmol) and (1-bromoethyl)benzene (0.124 mL, 0.909 mmol) were added to a Schlenk flask under continuous argon flux (positive pressure). The mixture was degassed by four freeze-pump-thaw cycles. After each thaw, the Schlenk was opened under argon flux (5 seconds), to remove the released gases from the solution and again closed under argon atmosphere. Then, CuBr (0.130 g, 0.909 mmol) and 5,5′-dimethyl-2,2′-dipyridyl (0.418 g, 2.27 mmol) were introduced into the frozen mixture under argon flux. The mixture was degassed by further two freeze-pump-thaw cycles. The Schlenk flask was then immersed in a pre-heated oil bath at 110 °C. After 18 h, THF (20.0 mL) was added to solubilize the solid product. The resulting mixture was permeated through a column filled with alumina (0.15 m length), using THF as an eluent in order to remove the metal complex catalyst from the polymeric solution. The polymeric solution was precipitated in methanol followed by drying under vacuum at 40 °C until reaching constant mass, thus obtaining a white solid product (PS-Br).

2.4 Preparation of amphiphilic PS-b-PVP copolymer via RAFT Macro-CTA was prepared by ionic substitution reaction, similar to a previously reported procedure[20,27]. In a two neck round bottom flask, potassium ethyl xanthogenate (0.267 g, 1.66 mmol) was dissolved in acetone (6.00 mL). Then, PS-Br (1.00 g) was dissolved in THF (5.00 mL) and added dropwise in the two neck round bottom flask under argon flux. This solvent mixture can be considered a good solvent for the components. The reaction was conducted at room temperature (25 ± 5 °C) overnight. To obtain the solid macro-CTA (PS-xanthogenate), the solution was precipitated in methanol followed by drying under vacuum at 40 °C until constant mass was obtained. The synthesis of the amphiphilic PS-b-PVP copolymer followed the traditional bulk RAFT polymerization. Macro‑CTA (1.0 g) and VP monomer (15.0 mL) were added in a Schlenk flask under continuous argon flow. The mixture was degassed by four freeze-pump-thaw cycles. Then, AIBN (0.198 mL, 0.0289 mmol) was introduced into the Schlenk flask under continuous argon flux. The mixture was degassed by further two freeze-pump-thaw cycles. The Schlenk flask was then immersed in a pre-heated oil bath at 70 °C for 48 h. Then, the viscous solution was precipitated in water, followed by centrifugation (13,500 rpm for 10 min) and drying under vacuum at 40 °C until constant mass was reached. Thus, the white solid product was purified by solubilization in THF, precipitation in distilled water, centrifugation and drying under the same conditions cited before.

2.5 Synthesis of silver nanoparticles using a PS-b-PVP copolymer A PS-b-PVP copolymer was used as a stabilizing agent in silver nanoparticle (AgNP) synthesis. In a typical experiment, 50 mg of PS-b-PVP copolymer were added to 10 mL of dimethylacetamide (DMAc) under stirring. Afterwards, 0.1 mL of aqueous NaOH (0.1 mol L-1) and 0.1 mL of aqueous AgNO3 (0.1 mol L-1) solutions were added drop-wise in 5 minute intervals. The colloidal suspension remained under magnetic stirring for a further 20 minutes. Polímeros, 26(1), 1-10, 2016


Synthesis and applications of polystyrene-block-poly(N-vinyl-2-pyrrolidone) copolymers 2.6 Synthesis of silica nanoparticles and PS-b-PVP copolymer adsorption Silica nanoparticles (SiO2NP) were prepared by the classical Stöber method[29]. Briefly, 2 mL of TEOS were added to 25 mL of ethanol within screw-cap vials, in the presence of 10.5 mmol of ammonium hydroxide solution. The alcoholic solution was kept under ultrasonic bath (25 kHz-200W) for 120 min at 30 °C[30]. The solid content was determined gravimetrically by drying the dispersion at 80 °C until a constant weight was obtained. The modification of SiO2NP was carried out as follows: a SiO2NP dispersion (1) was prepared by the addition of SiO2NP (50 mg) to a flask containing DMAc (10 mL) in an ultrasonic bath; a PS-b-PVP solution (2) was prepared by the dissolution of PS-b-PVP copolymers (50 mg) in DMAc (5 mL). Then, the solution (2) was slowly added to the dispersion (SiO2NP/DMAc) (1) under ultrasonic bath. After 48 h, the resulting dispersion was centrifuged at 14,000 rpm for 20 minutes. The supernatant containing excess of PS-b-PVP copolymer was removed and the precipitated solid (SiO2NP coated by PS-b-PVP copolymer) was lyophilized.

2.7 Characterization 1 H NMR spectra of the polymers were obtained in deuterated chloroform solutions (CDCl3) using a Bruker Avance III - 600MHz spectrometer operating at 600 MHz. Gel permeation chromatography (GPC) measurements were performed using a Viscotek GPCmax VE 2001, equipped with Viscotek VE 3580 RI Detector and three Shodex KF‑8060M columns working at 40 °C. Degassed THF was used as an eluent (1 mL min-1) and as a solvent to prepare 8.0 mg mL-1 sample solutions. Homopolymer and block copolymer molecular weights were determined using polystyrene standards. DSC analyses were performed with TA instruments DSC-Q100 in nitrogen atmosphere. Samples of ca. 5 mg were heated from 25 to 200 °C, cooled from 200 to -10 °C and heated from -10 to 200 °C at a heating rates of 10 °C min-1. TEM images were obtained using a Carl Zeiss LIBRA 120 PLUS operating at 120 kV and equipped with in-column OMEGA energy filter (EF-TEM). Samples were prepared by dropping 20 µL of a DMAc dispersion containing AgNP or SiO2NP in copper grid coated with thin amorphous carbon. Electron spectroscopic imaging (ESI‑TEM) was carried out for carbon mapping. The absorption spectra of colloidal Ag were obtained with an Agilent Cary 50 Probe ultraviolet–visible spectrophotometer using 0.5 mL of colloidal suspension diluted to 2.5 mL of DMAc.

3. Results and Discussion 3.1 Amphiphilic PS-b-PVP copolymers The synthesis of PS-Br homopolymers via ATRP is a well-established methodology in literature. According to 1 H NMR measurements (not shown here) the chain end functionality of the homopolymers synthesized was confirmed by the signals at 4.60-4.35 ppm attributed to the hydrogen located in the α position of the bromine chain end. Based on the signal intensity ratios between the 7.37-6.21 ppm range and those in the 4.60-4.35 ppm range, the calculated molecular weights by 1H NMR (MNMR), considering the ratio monomer:bromine chain end, were ca. 9,500 g mol-1 and 12,700 g mol-1, corresponding respectively to the PS91-Br and PS122-Br copolymers. Further Mn and Ɖ values obtained by GPC for the homopolymers and block copolymers synthesized, are shown in Table 1. The MNMR value represents the ratio of monomer units per bromine chain end. Depending on the application or the properties of interest, the use of the ratio between each block in the final amphiphilic block copolymer (e.g. 2:1 hydrophobic:hydrophilic block ratio) instead of the characteristic size of this block copolymer is more appropriate. For this reason, in the present work, we adopt the MNMR value because the resulting ratios of each block better represents the structural characteristics of the amphiphilic PS-b-PVP copolymer, since applications as coupling agent, surfactants, emulsifier or self-assembly micelles in selective solvent solutions are considered. The bromide PS chain end group was converted to xanthogenate by an ionic substitution reaction, producing a suitable macro chain transfer agent (macro-CTA) for the PS-b-PVP polymerization. Hussain et al.[20] produced amphiphilic PS-b-PVP copolymer starting from PS obtained by ATRP followed by RAFT of PVP. The authors used acetone as the solvent for conversion of the ATRP product (PS-Br) into RAFT macro-CTA, however acetone is not a solvent for this homopolymer[31]. Thus, in the present work, we used THF as it is a good solvent for polystyrene conversion into macro-CTA, avoiding phase separation and consequently low conversion yields due the higher molar masses of PS presented here. Similar conversion of the ATRP initiators into their corresponding RAFT macro-CTA was demonstrated by Wager et al.[27] applying modified ATRP conditions for polymer chain activation in presence of bis(thiobenzoyl) disulphide. Thus, the bromine chain end of poly(methylmethacrylate), poly(N,N-dimethylaminoethyl methacrylate) and poly(ethylene glycol) produced by ATRP were converted with high yields into to the desired RAFT end-functionality.

Table 1. Molecular parameters obtained for the homopolymers and block copolymers. Polymer

MNMR (g mol-1)

Mna (g mol-1)

Ɖa

Mass Yield (%) b

PS mass fraction c

9,500 12,700 17,600 20,500

5,800 9,600 10,500 13,900

1.5 1.2 1.5 1.4

73 68 65 69

0.55 0.64

PS91-Br PS122-Br PS91-b-PVP73 PS122-b-PVP70

Determined by GPC with THF as the eluente with respect to polystyrene standards. bCalculated after purification. cDetermined by 1H NMR.

a

Polímeros, 26(1), 1-10, 2016

3


de Farias, M. A., & Gonçalves, M. C. In this work, a xanthate containing an alkali metal (potassium) as R group was used to react with the PS homopolymer containing a halogen (bromine) as the chain end group, synthesized by ATRP. An ionic substitution reaction is easily carried out with these groups using adequate solvent for each species or a combination of solvents as used here. Based on this, we believe that the methodology used herewith can be applied to a vast range of other combinations of xanthates containing alkali metals as the R group and homopolymers synthesized by ATRP, thus leading to a great variety of functional macro-CTA appropriate to the RAFT technique. The key role of the RAFT technique is the appropriate choice and synthesis of the so-called CTA or macro-CTA[7,8], thus in this step, the PS xanthogenate chain end was used for the subsequent chain extension with the VP monomer in absence of solvent. The amphiphilic PS-b-PVP copolymer was initially synthesized via ATRP followed by RAFT for two reasons: (i) styrene is easily synthesized by ATRP technique, and in this case, the polymerization was conducted in bulk (solvent free); (ii) using only RAFT to synthesize PS homopolymer and subsequent PS-b-PVP copolymer would surely cause solubility problems due to the copolymer amphiphilic character. Furthermore, one could also point out the use of non-environmentally friendly solvents. Figure 1 shows the 1H NMR spectra of the synthesized amphiphilic PS-b-PVP copolymers. The PS hydrogen signals were assigned with “b”, “c”, “d”, “e” and “f”, xanthate hydrogen signals were assigned with “g” and “h”, and the PVP hydrogen signals were assigned with “i”, “j”, “k”, “l”, “o” and “p”. These assignments are also in agreement with literature[20,21,23,32]. Based on the signal intensity ratios between the 7.37–6.21 ppm range and those attributed only to PVP block (3.30 ppm and 3.90-3.55 ppm range), both copolymers were made up of around 70 VP monomeric units. The calculated MNMR of the block copolymers were ca. 17,600 g mol-1 and 20,500 g mol-1 corresponding to PS91-b-PVP73 and PS122-b‑PVP70 copolymers, respectively. Regarding the macro-CTA with the lowest molecular weight (PS91-xanthogenate), the PS91-b-PVP73 copolymer molecular weight almost doubled compared to the respective

homopolymer, corresponding to 0.55 PS mass fraction. Furthermore, no difference in Ɖ was observed, which remained the same for the homo and block copolymers. On the other hand, in relation to the macro-CTA with the highest molecular weight (PS122-xanthogenate), a slight increase in Ɖ was observed from the 1.2 to 1.4 and the corresponding block copolymer (PS122-b-PVP70) exhibited a 0.64 PS mass fraction. GPC curves of homopolymers and the respective block copolymer are shown in Figure 2. The GPC results of PS‑Br and PS-b-PVP copolymers displayed unimodal curves. The significant shift in these curves indicates a successful increase in the block copolymer molecular weights in relation to the original homopolymers. Other copolymers reported in literature, where the VP monomer was polymerized via RAFT from the PS macro-CTA, show low yields[33] or similar Ɖ values[21]. Bilalis et al.[23] synthesized VP block copolymers using styrene and 2-vinylpyridine via NMP and RAFT techniques, respectively. Applying the NMP technique, yields ranging from 20 to 70% were obtained, however a 2.2 Ɖ value was observed for the block copolymer with the highest yield. RAFT technique also provided similar results in relation to NMP, achieving yields ranging from 10 to 65% and 2.1 Ɖ value. The lowest Ɖ value (Ɖ = 1.4) was reached when PS was used as the first polymerized block followed by diblock synthesis with VP monomer addition. This value is in agreement with the results presented in Table 1. GPC measurements indicated that PS91-Br and PS122‑Br presented 1.5 and 1.2 Ɖ values, respectively (Table 1). One of the literature criteria used to define a controlled polymerization is based on the Ɖ values, where Mw/Mn ≤ 1.5 is considered a narrow molecular weight distribution[4,34-36]. Normally, very narrow Ɖ values (Mw/Mn ≤ 1.2) are obtained at low monomer conversion rates and/or by using solvent in the polymerization[25,37,38]. In the present work, we conducted the PS polymerization in bulk resulting Ɖ values slightly above very narrow Ɖ values (Mw/Mn ≤ 1.2) but still in the range of CRP Ɖ values (Mw/Mn ≤ 1.5).Thus, this work is in agreement with literature reported Ɖ values[4,6,20-22,25,35]. DSC heating curves of the PS-b-PVP copolymers is shown in Figure 3. Only the second heating scan curves

Figure 1. 1H NMR spectra of amphiphilic PS122-b-PVP70 copolymer and the corresponding peak assignments. 4

Polímeros, 26(1), 1-10, 2016


Synthesis and applications of polystyrene-block-poly(N-vinyl-2-pyrrolidone) copolymers

Figure 2. Comparison of GPC curves of (a) the lower and (b) the higher homopolymers with the corresponding block copolymers.

copolymers showed two distinct TG values, which are close to the values of the corresponding PS and PVP homopolymers, indicating that PS and PVP blocks are immiscible.

3.2 Use of PS-b-PVP as a AgNP stabilizing agent

Figure 3. DSC curves of the amphiphilic PS-b-PVP copolymers. The arrows show the Tg values of PS and PVP blocks.

are shown due to fact that they establish a defined sample history. Initially, the PS homopolymers presented glass transition temperatures (TG) of 97 and 100 °C for PS91-Br and PS122‑Br, respectively. On the other hand, with the formation of PS‑b‑PVP copolymers, TG values for the PS block increased about 5 °C in relation to the respective homopolymers, due to the increase in the molecular weight as well as the presence of the second rigid block. In the case of the PVP block, the TG values obtained for both copolymers was around 167 °C. Furthermore, in Figure 3, the transition around 167 °C is more defined for the PS91-b‑PVP73 than for the PS122-b-PVP70 copolymer. This result is in agreement with PVP ratio present in each block copolymer, since the PS91-b-PVP73 copolymer presented a higher PVP content and the transition around 167 °C was consequently easier to be detected for this block copolymer. Therefore, the PS-b-PVP Polímeros, 26(1), 1-10, 2016

According to the previous section, the produced block copolymers presented hydrophobic to hydrophilic ratio suitable to act as a stabilizing agent (hydrophobic:hydrophilic block ratios were approximately 1:1 and 2:1). PVP is a well-known polymer used as a stabilizing agent to produce metallic (e.g. silver, gold, etc) nanoparticles[39]. In the case of amphiphilic PS-b-PVP copolymer, PVP block can be used as stabilizing and co-reducing agents while the PS block can act as a hydrophobic segment. In this case, the amphiphilic PS-b-PVP copolymer can be used as a compatibilizer agent for polymeric matrices, mainly to those containing PS or a polymer miscible with PS. The synthesis of AgNP stabilized by PS-b-PVP (AgNP/PS-b‑PVP) was carried out by the direct addition of AgNO3 into the system (DMAc, NaOH and block copolymer). After AgNO3 addition, the solution turned yellowish indicating the successful formation of AgNP/PS-b-PVP in a unique step. In the present work, the PVP block size of both PS91‑b‑PVP73 and PS122-b-PVP70 copolymers synthesized are around 8,000 g mol-1, thus AgNP can be produced with any of the block copolymer synthesized in this work. Park et al.[19] synthesized the PS-b-PVP block copolymer (Mw = 20,000 g mol-1, PS mass fraction of 0.70 and Ɖ = 1.4) by ATRP using anisol as a solvent. The authors prepared a polymer solution dissolving PS-b-PVP (1 wt%) in THF. AgCF3CO3 (20 wt% with respect to block copolymer) was solubilized in the polymer solution and then the polymer/salt solution was dropped and spread on a glass slide in order to form a film. The glass slide containing the PS‑b‑PVP/ AgCF3CO3 film was UV irradiated (254 nm) for 1 h when the authors observed the AgNP/PS-b-PVP formation. The nanoparticles presented 4-6 nm average diameter, although UV-vis showed a broad band with low intensity. In this above related work, the AgNP were produced from two different steps: (i) PS-b-PVP/AgCF3CO3 film formation and (ii) posterior UV film irradiation. The PS-b-PVP synthesized by Park et al.[19] is very similar to PS122-b-PVP70 5


de Farias, M. A., & Gonçalves, M. C. synthesized in the present work (Mw = 19,000 g mol-1, PS mass fraction of 0.64 and Ɖ = 1.4). However, in this work, a one-step methodology was used to produce AgNP stabilized by PS‑b-PVP copolymers. The ultraviolet-visible (UV-vis) spectrum was used to confirm the formation of AgNP/PS-b-PVP. Only one absorption band with a 410 nm maximum absorbance was observed in the UV-vis spectrum presented in Figure 4a, confirming the characteristic silver surface plasmon[40,41] as well as the fact that the nanoparticles were spherical. The dispersity of colloidal suspension was evaluated by the full width at half maximum (FWHM) and the obtained value was 62 nm, which indicates that the AgNP/PS-b‑PVP can be considered monodispersed[42]. Typical AgNP/PS-b‑PVP EF‑TEM image are shown in Figure 4b, which present spherical monodisperse nanoparticles. These findings corroborate the UV-vis results. Furthermore, the AgNP/PS‑b-PVP diameter were measured from EF-TEM images and size distribution is shown in Figue 4 (c). The nanoparticle size distribution presented an 11 ± 1 nm average diameter. The AgNP/PS122-b-PVP70 were dried by lyophilization and submitted to a simple test: the AgNP/PS122-b-PVP70 was added to ethanol and sonicated for 10 minutes. After sonification, no AgNP redispersion was observed as shown in Figure 5a. However, removing ethanol and adding DMAc to the same container, the AgNP redispersion occurred immediately, as shown in Figure 5 (b). This simple test macroscopically demonstrates that the AgNP have an hydrophobic shell (PS block), which differentiates the AgNP produced in this work from the classical AgNP stabilized by PVP[11,43,44]. Thus, we can propose that there is a polymeric bilayer encapsulating the AgNP where the PVP block forms the inner layer and is responsible for the silver nanoparticle

stabilization, while the PS block forms the outer layer which provides a hydrophobic character to the nanoparticle external surface. For this reason, AgNP stabilized by PS-b-PVP copolymers do not redisperse in ethanol. The AgNP/PS-b-PVP colloidal suspensions were maintained unstirred in the dark for several weeks. No agglomeration was observed after 60 days, indicating that the PS-b-PVP copolymers provided stable AgNP suspensions.

3.3 Use of PS-b-PVP copolymer as surface modifier agent of SiO2NP Another application for the synthesized PS-b-PVP copolymer is a surface modifying agent. SiO2NP prepared by the Stöber method[29] was used to demonstrate this property. Nanoparticles prepared by this method presented diameters around 72 ± 9 nm. For this test, the PS-b-PVP solution was directly added to the SiO2NP/DMAc dispersions to promote copolymer adsorption on the nanoparticle surface. The PVP adsorption onto silica surface is attributed to the strong and spontaneous interaction between the carbonyl groups present in the amide ring and the silanol groups of the silica surface[45,46]. In relation to the adsorption of the PS-b-PVP copolymer on the SiO2NP surface, the formation of a polymeric bilayer encapsulating the nanoparticle can be proposed: the PVP block adsorbs preferentially at the SiO2NP surface, forming an inner layer, and the PS block forms an outer layer, which provides the SiO2NP with a hydrophobic surface character. Figure 6a and b compare EF-TEM images and elemental carbon maps of SiO2NP and SiO2NP modified by PS122‑b‑PVP70 (SiO2NP/PS122-b-PVP70), respectively. The green region in carbon maps corresponds to the carbon presence in the

Figure 4. (a) UV-vis spectra, (b) EF-TEM image and (c) size distribution of the Ag/PS-b-PVP. 6

Polímeros, 26(1), 1-10, 2016


Synthesis and applications of polystyrene-block-poly(N-vinyl-2-pyrrolidone) copolymers

Figure 5. (a) AgNP/PS122-b-PVP70 in ethanol and (b) after removing ethanol and adding DMAc. Insert in (a) shows the AgNP/PS122‑b‑PVP70 at the flask’sbottom after sonification.

Figure 6. EF-TEM images of SiO2NP synthesized by the Stöber method. (a) pure SiO2NP and (b) SiO2NP/PS122-b-PVP70 with their respective carbon maps.

respective EF-TEM bright field images. The SiO2NP carbon mapping exhibits the presence of carbon throughout the nanoparticles presented in Figure 6a. This mapping shows that the carbon content dispersed throughout the SiO2NP was, in fact, very small and is attributed to residual non-hydrolyzed ethoxy groups from the Stöber method synthesis[30,47,48]. Van Helden et al.[49] synthesized SiO2NP by the Stöber method and determined that the amount of non-hydrolyzed ethoxy groups was around 1 wt%. Polímeros, 26(1), 1-10, 2016

The PS-b-PVP copolymer layer adsorbed onto the SiO2NP nanoparticles can be seen in the EF-TEM image, Figure 6b. Furthermore, the corresponding carbon map of this image clearly showed the presence of a significant amount of carbon surrounding the nanoparticles (green regions), confirming the SiO2NP surface modification by the PS-b-PVP copolymer. The estimated average copolymer layer thickness was 5 ± 1 nm. Freris et al.[50] obtained similar results using poly(methyl methacrylate) to encapsulate 7


de Farias, M. A., & Gonçalves, M. C. SiO2NP with 180 nm diameter. In the related work, the authors identified a polymeric layer with thickness between 2 and 10 nm by TEM.

4. Conclusions PS-b-PVP copolymers with controlled molecular weight and relatively narrow polydispersity were synthesized using first ATRP for styrene polymerization and subsequently RAFT for N-vinyl-2-pyrrolidone polymerization. The PS-b-PVP copolymers presented number average molecular weight of 10,000-14,000 g mol-1, Mw/Mn ≤ 1.4 and 0.55-0.64 PS mass fraction. These PS-b-PVP copolymers were successfully applied as a stabilizing agent to synthesize AgNP in a single step, as confirmed by the UV-vis spectrum and EF-TEM images. No agglomeration was observed after 60 days, indicating that PS-b-PVP provides stable suspensions of AgNP in the organic solvent (dimethylacetamide). Furthermore, the PS-b-PVP copolymer also acted successfully as a surface modifying agent for the silica nanoparticles obtained by the Stöber method, as shown by EF-TEM images and elemental carbon maps. This surface modification provided SiO2NP with a hydrophobic surface character. This work highlighted the chemical and structural importance of PS‑b‑PVP copolymers in polymer science applications. Besides the applications shown here, these copolymers could also be used as water-in-oil emulsifier or coupling agents between inorganic particles and polymeric matrices.

5. Acknowledgements The authors thank PETROBRAS-CENPES, Brazilian agency FAPESP (FAPESP, 2010/17804-7) and National Institute of Science, Technology and Innovation in Complex Functional Materials (Inomat/INCT, FAPESP 2008/578678, CNPq 573644/2008-0) for the financial support. M. A. de Farias also thanks CAPES for the scholarship.

6. References 1. Wang, Z., Zhang, Q., Zhan, X., Chen, F., Rao, G., & Xiong, J. (2013). Preparation, kinetics and microstructures of well-defined PS-b-PS/Bd diblock copolymers via RAFT miniemulsion polymerization. Journal of Polymer Research, 20(11), 288. http://dx.doi.org/10.1007/s10965-013-0288-0. 2. Li, G. H., Yang, P. P., Gao, Z. S., & Zhu, Y. Q. (2012). Synthesis and micellar behavior of poly(acrylic acid-b-styrene) block copolymers. Colloid & Polymer Science, 290(17), 1825-1831. http://dx.doi.org/10.1007/s00396-012-2799-3. 3. Roghani-Mamaqani, H., Haddadi-Asl, V., Khezri, K., Zeinali, E., & Salami-Kalajahi, M. (2013). In situ atom transfer radical polymerization of styrene to in-plane functionalize graphene nanolayers: grafting through hydroxyl groups. Journal of Polymer Research, 21(1), 333. http://dx.doi.org/10.1007/ s10965-013-0333-z. 4. Wang, J.-S., & Matyjaszewski, K. (1995). Controlled/“living” radical polymerization. atom transfer radical polymerization in the presence of transition-metal complexes. Journal of the American Chemical Society, 117(20), 5614-5615. http://dx.doi. org/10.1021/ja00125a035. 5. Matyjaszewski, K., Patten, T. E., & Xia, J. (1997). Controlled/“living” radical polymerization. kinetics of the homogeneous atom transfer radical polymerization of styrene. 8

Journal of the American Chemical Society, 119(4), 674-680. http://dx.doi.org/10.1021/ja963361g. 6. Aitchison, T. J., Ginic-Markovic, M., Clarke, S., & Valiyaveettil, S. (2012). Polystyrene-block-poly(methyl methacrylate): initiation issues with block copolymer formation using ARGET ATRP. Macromolecular Chemistry and Physics, 213(1), 79-86. http://dx.doi.org/10.1002/macp.201100478. 7. Moad, G., Rizzardo, E., & Thang, S. H. (2005). Living radical polymerization by the RAFT process. Australian Journal of Chemistry, 58(6), 379-410. http://dx.doi.org/10.1071/CH05072. 8. Lowe, A. B., & Mccormick, C. L. (2007). Reversible addition – fragmentation chain transfer (RAFT) radical polymerization and the synthesis of water-soluble (co)polymers under homogeneous conditions in organic and aqueous media. Polymer, 32, 283351. http://dx.doi.org/10.1016/j.progpolymsci.2006.11.003. 9. Liu, X., Wu, Z., Zhou, F., Li, D., & Chen, H. (2010). Poly(vinylpyrrolidone-b-styrene) block copolymers tethered surfaces for protein adsorption and cell adhesion regulation. Colloids and Surfaces. B, Biointerfaces, 79(2), 452-459. http:// dx.doi.org/10.1016/j.colsurfb.2010.05.011. PMid:20554165. 10. Kumar, S., Changez, M., Murthy, C. N., Yamago, S., & Lee, J.-S. (2011). Synthesis of well-defined amphiphilic block copolymers by organotellurium-mediated living radical polymerization (TERP). Macromolecular Rapid Communications, 32(19), 1576-1582. http://dx.doi.org/10.1002/marc.201100277. PMid:21793088. 11. Zhang, Z., Zhao, B., & Hu, L. (1996). PVP protective mechanism of ultrafine silver powder synthesized by chemical reduction processes. Journal of Solid State Chemistry, 121(1), 105-110. http://dx.doi.org/10.1006/jssc.1996.0015. 12. Riess, G. (2003). Micellization of block copolymers. Progress in Polymer Science, 28(7), 1107-1170. http://dx.doi.org/10.1016/ S0079-6700(03)00015-7. 13. Hamley, I. W. (1998). The physics of block copolymers. Oxford: Oxford University Press. 14. Xue, B., Gao, L., Hou, Y., Liu, Z., & Jiang, L. (2013). Temperature controlled water/oil wettability of a surface fabricated by a block copolymer: application as a dual water/ oil on-off switch. Advanced Materials, 25(2), 273-277. http:// dx.doi.org/10.1002/adma.201202799. PMid:23074035. 15. Uyen, N. T. N., Joo, S. I., Kim, W. H., Oh, M. H., Lee, J., Lim, B. S., & Hong, S. C. (2013). Application of block copolymeric surface modifier with crosslinkable units for montmorillonite nanocomposites. Journal of Applied Polymer Science, 127(1), 690-698. http://dx.doi.org/10.1002/app.37856. 16. Zhang, D., Qi, L., Ma, J., & Cheng, H. (2001). Formation of silver nanowires in aqueous solutions of a double-hydrophilic block copolymer. Chemistry of Materials, 13(9), 2753-2755. http://dx.doi.org/10.1021/cm0105007. 17. Li, N., Qi, L., Shen, Y., Li, Y., & Chen, Y. (2013). Amphiphilic block copolymer modified magnetic nanoparticles for microwaveassisted extraction of polycyclic aromatic hydrocarbons in environmental water. Journal of Chromatography. A, 1316, 1-7. http://dx.doi.org/10.1016/j.chroma.2013.09.030. PMid:24119754. 18. Shamenkova, O. A., Mokeeva, L. K., Kopylova, N. A., & Semchikov, Y. D. (2006). Synthesis of amphiphilic block copolymers polystyrene-block-polyvinylpyrrolidone from active polystyrene. Russian Journal of Applied Chemistry, 79(3), 448-452. http://dx.doi.org/10.1134/S1070427206030232. 19. Park, J. T., Koh, J. H., Lee, K. J., Seo, J. A., Min, B. R., & Kim, J. H. (2008). Formation of silver nanoparticles created in situ in an amphiphilic block copolymer film. Journal of Applied Polymer Science, 110(4), 2352-2357. http://dx.doi. org/10.1002/app.28261. 20. Hussain, H., Tan, B. H., Gudipati, C. S., Liu, Y., He, C. B., & Davis, T. P. (2008). Synthesis and self-assembly of poly(styrene)Polímeros, 26(1), 1-10, 2016


Synthesis and applications of polystyrene-block-poly(N-vinyl-2-pyrrolidone) copolymers b-poly(N-vinylpyrrolidone) amphiphilic diblock copolymers made via a combined ATRP and MADIX approach. Journal of Polymer Science. Part A, Polymer Chemistry, 46(16), 56045615. http://dx.doi.org/10.1002/pola.22882. 21. Hu, D., & Zheng, S. (2010). Reaction-induced microphase separation in polybenzoxazine thermosets containing poly(Nvinyl pyrrolidone)-block-polystyrene diblock copolymer. Polymer, 51(26), 6346-6354. http://dx.doi.org/10.1016/j. polymer.2010.10.047. 22. Huang, C.-F., Nicolaÿ, R., Kwak, Y., Chang, F.-C., & Matyjaszewski, K. (2009). Homopolymerization and block copolymerization of N-vinylpyrrolidone by ATRP and RAFT with haloxanthate inifers. Macromolecules, 42(21), 8198-8210. http://dx.doi.org/10.1021/ma901578z. 23. Bilalis, P., Pitsikalis, M., & Hadjichristidis, N. (2006). Controlled nitroxide-mediated and reversible addition-fragmentation chain transfer polymerization of N-vinylpyrrolidone: Synthesis of block copolymers with styrene and 2-vinylpyridine. Journal of Polymer Science. Part A, Polymer Chemistry, 44(1), 659-665. http://dx.doi.org/10.1002/pola.21198. 24. Arsalani, N., Fattahi, H., & Entezami, A. A. (2006). Synthesis of amphiphilic diblock and random copolymers of styrene and N-vinylpyrrolidone using nitroxide-mediated living free radical polymerization. Iranian Polymer Journal, 15(12), 997-1005. Retrieved in 16 January 2015, from http://www. sid.ir/en/vewssid/j_pdf/81320061207.pdf 25. Ray, B., Kotani, M., & Yamago, S. (2006). Highly controlled synthesis of poly(N-vinylpyrrolidone) and its block copolymers by organostibine-mediated living radical polymerization. Macromolecules, 39(16), 5259-5265. http://dx.doi.org/10.1021/ ma060248u. 26. Zakharova, O. G., Golyagina, Y. V., & Semchikov, Y. D. (2009). Synthesis and surface properties of amphiphilic block copolymers polyvinylpyrrolidone-block-polystyrene. Russian Journal of Applied Chemistry, 82(4), 644-649. http://dx.doi. org/10.1134/S107042720904020X. 27. Wager, C. M., Haddleton, D. M., & Bon, S. A. (2004). A simple method to convert atom transfer radical polymerization (ATRP) initiators into reversible addition fragmentation chain-transfer (RAFT) mediators. European Polymer Journal, 40(3), 641-645. http://dx.doi.org/10.1016/j.eurpolymj.2003.10.025. 28. Hayes, W., & Rannard, S. (2004). Controlled/“living” polymerization methods. In F. J. Daves (Ed.). Polymer synthesis – a practical approach (pp. 99-125). New York: Oxford University Press. 29. Stöber, W., Fink, A., & Bohn, E. (1968). Controlled growth of monodisperse silica spheres in the micron size range. Journal of Colloid and Interface Science, 26(1), 62-69. http://dx.doi. org/10.1016/0021-9797(68)90272-5. 30. Costa, C. A. R., Leite, C. A. P., & Galembeck, F. (2003). Size dependence of Stöber silica nanoparticle microchemistry. The Journal of Physical Chemistry B, 107(20), 4747-4755. http:// dx.doi.org/10.1021/jp027525t. 31. Mark, J. E. (1999). Polymer data handbook. New York: Oxford University Press. 32. Williams, D., & Fleming, I. (2007). Spectroscopic methods in organic chemistry (6th ed.). New York: McGraw-Hill Higher Education. 33. Wan, D., Satoh, K., Kamigaito, M., & Okamoto, Y. (2005). Xanthate-mediated radical polymerization of N-vinylpyrrolidone in fluoroalcohols for simultaneous control of molecular weight and tacticity. Macromolecules, 38(25), 10397-10405. http:// dx.doi.org/10.1021/ma0515230. 34. Wang, J.-S., & Matyjaszewski, K. (1995). Controlled/“living” radical polymerization. Halogen atom transfer radical polymerization promoted by a Cu(I)/Cu(II) redox process. Polímeros, 26(1), 1-10, 2016

Macromolecules, 28(23), 7901-7910. http://dx.doi.org/10.1021/ ma00127a042. 35. Matyjaszewski, K., & Xia, J. (2001). Atom transfer radical polymerization. Chemical Reviews, 101(9), 2921-2990. http:// dx.doi.org/10.1021/cr940534g. PMid:11749397. 36. Perrier, S., & Takolpuckdee, P. (2005). Macromolecular design via reversible addition-fragmentation chain transfer (RAFT)/ xanthates (MADIX) polymerization. Journal of Polymer Science. Part A, Polymer Chemistry, 43(22), 5347-5393. http://dx.doi. org/10.1002/pola.20986. 37. Kahveci, M. U., Acik, G., & Yagci, Y. (2012). Synthesis of block copolymers by combination of atom transfer radical polymerization and visible light-induced free radical promoted cationic polymerization. Macromolecular Rapid Communications, 33(4), 309-313. http://dx.doi.org/10.1002/marc.201100641. PMid:22253209. 38. Chavda, S., Yusa, S., Inoue, M., Abezgauz, L., Kesselman, E., Danino, D., & Bahadur, P. (2013). Synthesis of stimuli responsive PEG47–b-PAA126–b-PSt32 triblock copolymer and its self-assembly in aqueous solutions. European Polymer Journal, 49(1), 209-216. http://dx.doi.org/10.1016/j. eurpolymj.2012.09.021. 39. Wang, H., Qiao, X., Chen, J., Wang, X., & Ding, S. (2005). Mechanisms of PVP in the preparation of silver nanoparticles. Materials Chemistry and Physics, 94(2-3), 449-453. http:// dx.doi.org/10.1016/j.matchemphys.2005.05.005. 40. Noguez, C. (2007). Surface plasmons on metal nanoparticles: the influence of shape and physical environment. The Journal of Physical Chemistry C, 111(10), 3806-3819. http://dx.doi. org/10.1021/jp066539m. 41. Indumathy, R., Sreeram, K. J., Sriranjani, M., Aby, C. P., & Nair, B. U. (2010). Bifunctional role of thiosalicylic acid in the synthesis of silver nanoparticles. Materials Sciences and Applications, 1(05), 272-278. http://dx.doi.org/10.4236/ msa.2010.15040. 42. Medina-Ramirez, I., Bashir, S., Luo, Z., & Liu, J. L. (2009). Green synthesis and characterization of polymer-stabilized silver nanoparticles. Colloids and Surfaces. B, Biointerfaces, 73(2), 185-191. http://dx.doi.org/10.1016/j.colsurfb.2009.05.015. PMid:19539451. 43. Lee, J.-M., Jun, Y.-D., Kim, D.-W., Lee, Y.-H., & Oh, S.-G. (2009). Effects of PVP on the formation of silver–polystyrene heterogeneous nanocomposite particles in novel preparation route involving polyol process: Molecular weight and concentration of PVP. Materials Chemistry and Physics, 114(2-3), 549-555. http://dx.doi.org/10.1016/j.matchemphys.2008.10.001. 44. Wiley, B. J., Im, S. H., Li, Z.-Y., McLellan, J., Siekkinen, A., & Xia, Y. (2006). Maneuvering the surface plasmon resonance of silver nanostructures through shape-controlled synthesis. The Journal of Physical Chemistry B, 110(32), 15666-15675. http://dx.doi.org/10.1021/jp0608628. PMid:16898709. 45. Cohen Stuart, M., Fleer, G., & Bijsterbosch, B. (1982). The adsorption of poly(vinyl pyrrolidone) onto silica. I. Adsorbed amount. Journal of Colloid and Interface Science, 90(2), 310320. http://dx.doi.org/10.1016/0021-9797(82)90300-9. 46. Robinson, S., & Williams, P. A. (2002). Inhibition of protein adsorption onto silica by polyvinylpyrrolidone. Langmuir, 18(23), 8743-8748. http://dx.doi.org/10.1021/la020376l. 47. Costa, C. A. R., Leite, C. A. P., & Galembeck, F. (2006). ESITEM imaging of surfactants and ions sorbed in Stöber silica nanoparticles. Langmuir, 22(17), 7159-7166. http://dx.doi. org/10.1021/la060389p. PMid:16893211. 48. Costa, C. A. R., Valadares, L. F., & Galembeck, F. (2007). Stöber silica particle size effect on the hardness and brittleness of silica monoliths. Colloids and Surfaces. A, Physicochemical and Engineering Aspects, 302(1-3), 371-376. http://dx.doi. org/10.1016/j.colsurfa.2007.02.061. 9


de Farias, M. A., & Gonçalves, M. C. 49. Van Helden, A. K., Jansen, J. W., & Vrij, A. (1981). Preparation and characterization of spherical monodisperse silica dispersions in nonaqueous solvents. Journal of Colloid and Interface Science, 81(2), 354-368. http://dx.doi.org/10.1016/00219797(81)90417-3. 50. Freris, I., Cristofori, D., Riello, P., & Benedetti, A. (2009). Encapsulation of submicrometer-sized silica particles by a

10

thin shell of poly(methyl methacrylate). Journal of Colloid and Interface Science, 331(2), 351-355. http://dx.doi.org/10.1016/j. jcis.2008.11.052. PMid:19081575. Received: Jan. 16, 2015 Revised: Aug. 13, 2015 Accepted: Aug. 17, 2015

PolĂ­meros, 26(1), 1-10, 2016


http://dx.doi.org/10.1590/0104-1428.2041

Curing and thermal behavior of epoxy resins of hexafluoro - bisphenol –A and bisphenol-A Vaishnav Kiran1 and Bharti Gaur1* Research and Development Laboratory, Department of Chemistry, National Institute of Technology Hamirpur, Himachal Pradesh, India

1

*bhartigaur@gmail.com

Abstract This paper describes the synthesis and characterization of epoxy resins based on (hexafluoroisopropylidene)diphenol (EFN) and p,p’-isopropylidenebisphenol (EBN), respectively and 4, 4’- (hexafluoroisopropylidene)dipthalic-imideamine (IMAM), a curing agent. The synthesized epoxy resins and IMAM curing agent were characterized by Fourier Transform Infrared (FTIR) and 1H Nuclear Magnetic Resonance (NMR) spectroscopy.13C NMR technique was also used to characterize IMAM. Study of curing behavior of EFN and EBN with stoichiometric amount of aromatic 4,4’-diaminodiphenylmethane (DDM), 4,4’-diaminodiphenylsulfone (DDS) and IMAM by using Differential Scanning Calorimetery (DSC) indicated that IMAM was least reactive curing agent towards both epoxy resins as compared to DDS and DDM. The investigation of thermal decomposition of the cured compounds by thermogravimetric analyzer (TGA) indicated the higher thermal stability of EFN and EBN resins initially with DDS and at elevated temperatures with IMAM. It was also observed that EFN resins were thermally more stable than EBN resins cured with corresponding curing agents. Keywords: aromatic diamines, curing behavior, epoxy resins, imide amine, thermal properties.

1. Introduction Epoxy resins are one of the most important thermosetting resins known to possess good mechanical properties, high chemical resistance and excellent adhesive properties. These desirable properties make them widely applicable materials viz for surface coatings, adhesives, corrosion protectants, composites and laminates, encapsulants for semiconductors, insulating materials for electronic devices etc[1-10]. However the conventional epoxy resins do not exhibit very high thermal resistance, which is a prerequisite for advanced materials. Over the years multifunctional epoxy resins have been synthesized by many researchers to improve the thermal stability of these resins, either by incorporating aromatic ring as in case of phenolic based resins or introducing phosphorus in the epoxy backbone[11-13]. Attempts have also been made to introduce fluorinated substituents into the polymer backbone in order to improve electrical properties such as dielectric constant. These fluorinated substituents are also reported to reduce the moisture absorption due to the nonpolar character of fluorocarbons[14]. Some typical structures of multifunctional epoxy resin are shown in Figure 1. These resins are reported to be used as matrix material for high performance fiber-reinforced composites in the aerospace industry and as encapsulant for electronic components. Outstanding properties and performance of epoxy resins are obtained by crosslinking these into a three-dimensional, insoluble and infusible network by reacting with suitable curing agents. The crosslinking occurs through reaction of the epoxide or oxirane group. In the recent years attention has also been focused by the researchers on the development of novel curing agents that will help to improve the performance of epoxy resins at elevated temperatures. This article reports the introduction of pendant fluoro groups on the epoxy backbone in order to increase the thermal stability of the

Polímeros, 26(1), 11-20, 2016

epoxy resin. An attempt has also been made to host these fluorinated substituents as well as the aromatic rings on the backbone of the curing agent, an Imide-amine, in order to obtain matrix which could show even better heat stability. Curing and decomposition behavior of fluorine containing epoxy resin (EFN) was compared with bisphenol-A novolac based (EBN) epoxy resin by using aromatic DDM, DDS and IMAM as curing agents.

2. Materials and Methods 2.1 Materials 4,4’-diaminodiphenlymethane, 4,4’-diaminodiphenyl sulfone, p,p’-isopropylidenebisphenol (bisphenol-A/BPA) and 4,4’- (hexafluoroisopropylidene) dipthalicanhydride were purchased from Alfa Aesar. Formaldehyde, tetrahydrofuran (THF) and glacial acetic acid were received from Fisher Scientific, Epichlorohydrin (ECH) and sodium hydroxide were obtained from Loba Chemie. Isopropanol and methyl isobutyl ketone (MIBK) were purchased from Merck, 4,4’-(hexafluoroisopropylidene)diphenol (hexafluoro/6Fbisphenol-A/6F-BPA) was obtained from Sigma Aldrich and paratoluene sulfonic acid (PTSA) from laboratory reagents, sodium bicarbonate was of analytical grade.

2.2 Synthesis and Characterization 2.2.1 Synthesis of 4, 4’-(hexafluoro-isopropylidene) - diphenol (6F-BPA) and p,p’-isopropylidenebisphenol (BPA) based novolac resins (Scheme 1 and 2) A 40ml methyl isobutyl ketone (MIBK) solution containing 6F-BPA (0.05 mol) and PTSA (0.0014 mol) was added to a 250ml three- necked round bottom flask equipped with

11

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


Kiran, V., & Gaur, B.

Figure 1. Some typical structures of multifunctional epoxy resins[11-13].

Scheme 1. Synthesis of hexafluorobisphenol-A novolac epoxy resin (EFN).

mechanical stirrer and dean and stark trap with a reflux condenser. Nitrogen gas was purged for 30 minutes and the reaction solution was heated to 100°C with stirring. 0.08 mol of formaldehyde solution (37-41%) was added to the reaction mixture drop wise. Then, the reaction mixture was heated to 120°C with constant stirring and maintained at this temperature for five hours. The condensation of novolac 12

resin is a reversible process; therefore (0.13 mol) of water generated during the reaction was removed as azeotropic mixture with MIBK. The reaction mixture was cooled to room temperature and washed with deionised water several times until it became neutral. The solution was then distilled at 120°C in order to remove the MIBK solvent. The product was then washed with a mixture of water/methanol (2:1 v/v) Polímeros, 26(1), 11-20, 2016


Curing and thermal behavior of epoxy resins of hexafluoro - bisphenol –A and bisphenol-A several times to remove the unreacted 6F-BPA. A red brown solid product was obtained after vacuum drying at 60°C for 48 h.Similar procedure was followed for the synthesis of BPA based novolac resin. 2.2.2 Synthesis of novolac epoxy resins of 4, 4’-(hexafluoroisopropylidene) - diphenol (EFN) and p,p’- isopropylidenebisphenol (EBN) (Scheme 1 and 2) To the above dried product of 6F-BPA novolac 8 mol of epichlorohydrin for every phenolic group of novolac resin and isopropyl alcohol (0.83mol) was added to 250ml three necked round bottom flask equipped with mechanical stirrer and condenser. After increasing the reaction temperature to 70°C with constant stirring, 0.078 mol of 20 wt% aqueous solution of sodium hydroxide was added drop wise into the reaction solution within one hour. The system was maintained at 70°C for another four hours with constant stirring. The reaction product in the flask was washed several times with deionized water to remove the sodium chloride. The product was dissolved in toluene and filtered in order to remove residual salt. Toluene and excess of ECH were distilled off under reduced pressure. Finally the product obtained was dried at 60-70°C under vacuum for 48 hours. BPA based novolac epoxy resin (EBN) was also synthesized by adopting a similar procedure. 2.2.3 Synthesis of 4, 4’- (hexafluoroisopropylidene)diphthalicimideamine (IMAM) curing agent (Scheme 3) 0.011 mol of DDM was dissolved in 40 ml glacial acetic acid in a round bottom flask equipped with condenser. To this 0.005mol of 4, 4’- (hexafluoroisopropylidene)diphthalic anhydride was added. The reaction mixture was heated at

120°C for ten hours with constant stirring. The prepared 4, 4’-(hexafluoroisopropylidene)diphthalic imideamine (IMAM) was precipitated in ice-cold water, filtered and washed several times with distilled water followed by washings with sodium bicarbonate solution. The product was again washed with distilled water several times before drying it in vacuum oven at 50-60°C.

2.3 Characterization 2.3.1 Structural characterization FTIR spectra of the samples were recorded by using Perkin Elmer 1600 FTIR spectrophotometer in the range of 4000-400cm-1 on the KBr pellets. 1H NMR and 13C NMR spectra were recorded on a BRUKER AVANCE II 400 NMR spectrometer using deuterated Dimethylsulfoxide as solvent, and tetramethylsilane as the internal standard. 2.3.2 Curing and decomposition behavior The samples for curing studies were freshly prepared at 25°C by mixing the novolac epoxy resins EFN and EBN with three different curing agents DDM/DDS/IMAM, dissolved in appropriate solvents respectively, in the molar ratio of 1:2 in small glass vial with vigorous stirring using a glass rod until the samples became homogeneous in nature. DDM and DDS were dissolved in minimum amount of acetone whereas IMAM in minimum amount of THF. The solvents were then removed by vacuum stripping before recording the DSC scans. DSC scans for all these samples under dynamic conditions were obtained on Mettler Differential Scaning Calorimetery with programmed heating rate of 10°C min-1

Scheme 2. Synthesis of bisphenol-A novolac epoxy resin (EBN).

Scheme 3. Synthesis of 4, 4’- (hexafluoroisopropylidene)diphthalic-imideamine (IMAM). Polímeros, 26(1), 11-20, 2016

13


Kiran, V., & Gaur, B. from 35°C to the temperature at which exothermic reactions were completed. Thermal stability of the samples cured isothermally at 250±10 °C in hot air oven was evaluated by thermogravimetry in nitrogen atmosphere (flow rate=200ml/min). EXSTAR TG/DTA 6300 was used to record TG/DTG/DTA traces at heating rate of 10°C min-1 with sample size of 10±1 mg. The relative thermal stabilities of the different cured resins were quantitatively estimated by comparing the temperature at which maximum degree of weight loss occurred and by determining the activation energy of the thermal decomposition reaction.

2.3.3 Epoxide equivqlent weight Epoxide equivalent weights (EEWs) of synthesized epoxy resins EFN and EBN were determined by the pyridinium chloride method[15] were found to be 354 and 377 respectively.

3. Results and Discussions 3.1 Characterization of 6F-BPA and BPA novolacs and their epoxy resins The FTIR spectra of 6F-Bisphenol-A novolac (6F-BPA-N) (Figure 2a) and Bisphenol-A novolac (BPA-N) resins showed the –CH2- stretching vibration due to bridging

Figure 2. FTIR spectra of (a) 6F-BPA novolac; (b) EFN novolac epoxy resin. 14

Polímeros, 26(1), 11-20, 2016


Curing and thermal behavior of epoxy resins of hexafluoro - bisphenol –A and bisphenol-A methyene groups at 2965 cm-1 . A broad band was observed at 3353 cm-1due to the hydroxyl group. Apart from the other peaks appearing for the aromatic C-H stretching, C=C stretching, -CH2- bending, as in case of BPA-N resin. IR spectrum of 6F-BPA-N resin showed peaks in the region of 1372-1446 cm-1, which indicated the presence of C-F bonds. The FTIR spectra of EFN (Figure 2b) and EBN novolac epoxy resins showed characteristic absorptions at 3057 and 2929cm-1 due to streching vibrations of aromatic rings and bridging methylene groups, respectively. A characteristic absorption band at 1299 cm-1 depicted the ring breathing frequency of epoxy ring, the appearance of band at 916 cm-1 proved the asymmetric ring deformation and band at 762 cm-1 showed the symmetric ring deformation of epoxy ring. The spectrum of EFN (Figure 2b) also showed peaks in between 1372-1446 cm-1 due to the presence C-F bonds. Figure 3 shows the 1H NMR spectra of 6F-BPA-N, BPA-N novolac resins and EFN, EBN novolac epoxy resins. Figures 3a and c showed characteristic proton resonance signal at 3.7 ppm due to methylene (–CH2) bridging protons and multiplet for aromatic protons at 6.6-7.1 ppm. A singlet due to -OH protons of phenol appeared at 9.5 ppm in the spectrum of 6F-BPA-N (Figure 3a) and at 8.7 ppm in the spectrum of BPA-N (Figure 3c). BPA-N resin also showed a characteristic singlet due to –CH3 protons at 1.5 ppm (Figure 3c). 1HNMR spectra (Figures 3b and 3d) of EFN and EBN showed proton resonance signals at 3.6-3.8 and 6.7-7.3 ppm due to –CH2 bridging and aromatic protons respectively. Characteristic proton resonance signals at

2.7-2.9, 3.3 and 3.9-4.2 ppm due to –O-CH2, -CH and -CH2 protons of epoxy ring were also observed, respectively. 3.1.1 Characterization of IMAM curing agent The FTIR spectrum of IMAM curing agent in Figure 4a showed characteristic absorption bands at 1783 and 1725 cm-1 due to asymmetric and symmetric stretching vibration of imide group and a band was observed at 722 cm-1 due to five-membered ring deformation of the cyclic imide. –NH2 stretching bands appeared at 3260 ± 30 cm-1. Absorption bands for C-N and C=C stretching of aromatic rings were observed at 1409-1435 cm-1 and 1602-1672 cm-1, respectively.1H NMR spectrum Figure 4b of IMAM curing agent showed proton resonance signal at 3.9 ppm for –NH2 protons and 7.0-8.1 ppm for aromatic protons. A resonance signal appeared at 2.0 ppm due to methylene protons.13C NMR spectrum (Figure 4c) was recorded in order to confirm the cyclization of the amic acid which is an intermediate in the reaction represented by the Scheme 3. The characteristic signal at 168 ppm due to cyclic imide carbon was observed. Carbons of Ar-CH2 and Ar-NH2 appeared at 132 and 141 ppm respectively. Peaks due to aromatic carbons were observed at 119-137 ppm.

3.2 Thermal curing behavior The curing behavior of EFN and EBN novolac epoxy resins with three different curing agents were studied by DSC. The reactivity of the curing agents towards epoxy resins can be examined from the onset exothermic temperature. The onset exothermic temperature was obtained by extrapolating the steepest portion of the curve to the base line. A curing agent

Figure 3. 1HNMR of (a) 6F-BPA novolac; (b) EFN novolac epoxy resin; (c) BPA novolac; (d) EBN novolac epoxy resin. Polímeros, 26(1), 11-20, 2016

15


Kiran, V., & Gaur, B. that exhibits lower onset exothermic temperature is more reactive towards the epoxy resin. Figures 5 and 6 show the typical DSC scans for curing of EFN and EBN epoxies with DDS and IMAM curing agents respectively. The exothermic temperatures, onset, peak and endset temperatures have been given in Table 1. The enthalpy change (ΔH) which can be used to determine the extent of curing of resin[16,17] have been calculated by measuring the area under the

exothermic transition and are also summarized in Table 1. The onset curing exothermic temperature of EFN with DDM/DDS/IMAM were observed at 118, 184, 240°C, respectively. Curing exotherms of EBN with DDM/DDS/ IMAM showed the onset curing temperatures at 126, 147, 256°C respectively. The reaction of the amine with an oxirane ring is a nucleophilic addition reaction. From the results of thermal curing it was observed that despite the presence of

Figure 4. (a) FTIR; (b) 1H NMR of IMAM curing agent; (c) 13 C NMR of IMAM curing agent.

Figure 5. DSC scans of novolac based epoxy resins of (a) EFN/DDS; (b) EFN/IMAM. 16

Polímeros, 26(1), 11-20, 2016


Curing and thermal behavior of epoxy resins of hexafluoro - bisphenol –A and bisphenol-A

Figure 6. DSC scans of novolac based epoxy resins of (a) EBN/DDS; (b) EBN/IMAM. Table 1. DSC results of cured epoxy resins. Sample Designation

Tonset (°C)

Tpeak (°C)

Tendset (°C)

ΔH (J/g)

EFN- DDM EFN- DDS EFN-IMAM EBN- DDM EBN- DDS EBN- IMAM

118 184 240 126 147 256

144 235 273 134 210 293

194 280 295 154 259 315

99.81 65.07 56.14 73.88 64.22 48.30

electron donating methyl group in IMAM it was found to be the least reactive. The reason for this may be attributed to the presence of electron withdrawing CF3 and carbonyl groups in the main chain as well as the steric hinderance caused due to the bulkiness of the groups. This reduced the basicity of amine and consequently the reactivity of the IMAM curing agent decreased. Whereas in DDS only electronic factor dominated and owing to the electron withdrawing nature of sulfone group the reactivity of DDS was found to be higher than IMAM. The highest reactivity of DDM towards both the epoxy resins was in good agreement with electronic donating effects of methane moiety[1,10,18]. Thus the order of basicity of amine groups in the three curing agents will be DDM>DDS>IMAM. Furthermore, the melting point of DDS (175°C) and IMAM (195°C) is very high as compared to DDM (90°C), which may also have a negative effect on their reactivity. The EFN epoxy resin was observed to be more reactive towards the curing agents as compared to that EBN which may be probably due to electronic factors. The presence of electron withdrawing groups CF3 may facilitate the oxirane ring opening during the reaction with the amine group.

3.3 Thermal properties of cured resins EFN and EBN epoxy resin samples were isothermally cured at 250±10 °C for twelve hours with stoichiometric amount of diamines or imide-amine curing agents. DSC scans were recorded for these samples and absence of any residual exothermic peak in their scans confirmed complete curing of the samples. TGA technique considered to be the most Polímeros, 26(1), 11-20, 2016

favorable technique for rapid estimation in comparing and ranking the thermal stability of samples, was employed for determining the thermal stability of these isothermally cured samples. The relative thermal stability of the cured samples was compared by observing the initial degradation temperatures (IDT), temperature of maximum degradation (Tmax), final degradation temperature (FDT) and residual weight retention (Rw) at 900°C. The results have been summarized in Table 2. The TG curves of cured epoxy resins EFN/DDM, EFN/DDS, EFN/IMAM are shown in Figure 7. It can be seen that the samples did not show initial weight loss till 345°C which implied that the cured samples of EFN resin were thermally stable till this temperature. As the temperature reached 350±5°C these samples showed a slight weight loss and it was observed that the EFN/DDS sample showed better thermal stability as compared to EFN/DDM and was found to be thermally less stable than that of EFN/IMAM. However as the scanning temperature reached 415±5°C rapid thermal degradation occurred and EFN/DDS found to be less stable as compared to corresponding EFN/DDM and EFN/IMAM respectively and the final decomposition temperatures of all the samples were found to be almost identical. Figure 8a shows typical TG/DTG/DTA scans of cured epoxy resin of EBN with IMAM and Figure 8b shows the TG curves for cured EBN/DDM, EBN/DDS and EBN/IMAM, respectively. From the Figure 8b and Table 2 it can be observed that the cured EBN resins containing DDM/DDS/IMAM, respectively, were thermally stable till the scanning temperature reached up to 330±5°C. As the 17


Kiran, V., & Gaur, B. Table 2. Thermal properties of cured epoxy resins. Sample Designation EFN-DDM EFN-DDS EFN-IMAM EBN-DDM EBN-DDS EBN-IMAM

IDT (°C) 354 359 369 339 350 356

Tmax (°C)

FDT (°C)

422 417 456 409 407 453

540 542 544 465 470 490

Residual weight retention (Rw) % 26.46 22.48 29.08 22.83 17.65 30.80

Activation energy of decomposition reaction (Ea) Temperature Range (°C) 290-690 290-625 280-670 265-560 280-560 270-570

Ea (KJ/mol) 62.825 48.742 74.208 63.836 59.773 96.365

Figure 7. TG curves for EFN cured with DDM/DDS/IMAM.

scanning temperature rose above 340°C, the cured EBN/DDS sample showed the values of IDT, Tmax, FDT as 350, 407, and 470°C respectively. However the FDT for the sample EBN/IMAM was found to be higher as compared EBN cured with DDM and DDS respectively. From the values of IDT and FDT it was observed that initially better thermal stability of EFN and EBN resins with DDS may be due to the rigidity provided by high functionality and rigid structure of DDS as compared to that cured with DDM[19]. For the superior thermal stability of EFN and EBN resins cured with IMAM, It is assumed that the solid state aromatization reaction in the case of IMAM are promoted by the formation of the crosslinks which may develop between the carbonyl and trapped or unreacted amino groups at elevated temperatures and thus may have increased the thermal stability of the sample. It was also seen that the cured EFN and EBN resins with IMAM showed highest residual weight retentions (R w) at 900°C. This was probably due to the presence of thermally oxidative resistant CF 3 groups in the main chain of IMAM. Furthermore the cured EFN resins were found to be thermally more stable as compared to cured EBN resins. The enhanced thermal properties of cured EFN samples were probably due to the presence of fluorinated groups in the EFN network structure. 18

Figure 8. (a) Typical TG/DTG/DTA scans for cured sample of EBN/IMAM; (b) TG curves for EBN cured with DDM/DDS/ IMAM.

3.3.1 Activation energy of decomposition (Ea) Activation energy of decomposition (Ea) can also be used as a decisive factor for evaluating the thermal stability of the polymers, higher the value of activation energy higher will be the stability of the compounds. Because the rate constant of reaction depends upon energy difference between initial and transistion states of the reaction. If higher is the value of the energy difference higher will be the activation energy and more will be the stability of the compound. Calculation of activation energy from the TG curves has Polímeros, 26(1), 11-20, 2016


Curing and thermal behavior of epoxy resins of hexafluoro - bisphenol –A and bisphenol-A been reported in the literature by several methods[20-24]. But the expression used by Dharwadkar and Kharkhanawala for the calculation of activation energy is independent of sample size and heating rate. Ea100θ ln ln = +C (1 − a )−1   RTi2 (T f − Ti )

(1)

where α, fraction reacted; Ti, the temperature of inception of reaction; Tf, temperature at the point of inflection on the thermogram; θ = (T-Ts) (where T is the temperature under consideration and Ts is the maximum temperature); C, −1 constant; R, gas constant. A plot of ln ln (1 − a )  versus   θ gave a straight line was equal to m=

100 Ea

RTI 2 (T f −Ti )

(2)

From this equation, Ea was calculated.

Activation energies of decomposition (Ea) were calculated by using Equations 1 and 2 and the values of E have been summarized in Table 2. On comparing the activation energies of decomposition reaction of EFN and EBN novolac epoxy resins cured with DDM/ DDS/ IMAM curing agents respectively, the following trend was observed for activation energies [(EFN / EBN) + IMAM ] > [(EFN / EBN) + DDM ] > [(EFN / EBN) + DDS]. These results also substantiate the fact that the presence of fluorinated groups in the curing agent IMAM have also contributed towards the improvement of thermal stability of both the cured epoxy resins. From Table 2 it was also observed that the cured samples having high Rw value showed higher activation energy which led to the increase in their thermal stability.

4. Conclusions EFN and EBN novolac epoxy resins were synthesized, which were then thermally cured with aromatic amines (DDM, DDS) and synthesized aromatic imide-amine (IMAM) curing agents. The investigation of curing behavior of cured epoxy resins revealed that IMAM was found to be the least reactive curing agent towards both EFN and EBN as compared to their counterparts DDM and DDS. This was attributed to the electronic effects, high molecular weight and physical properties of IMAM curing agent. The experimental results for the thermal degradation demonstrated that the major weight loss for all the samples occurred in the temperature range of 300-600°C. Both the EFN and EBN samples cured with IMAM exhibited higher thermal stability probably due to the solid state aromatization reactions which may lead to the formation of cross-links between the carbonyl and trapped or unreacted amino groups. The presence of CF3 groups in IMAM may also have played its role in the enhancement of thermal stability of such cured resins. The EFN cured resins were found to be more heat resistant as compared to their corresponding EBN resins. This was owing to the ease of oxirane ring opening due to the presence of CF3 groups in EFN that might have been a good ignition for Polímeros, 26(1), 11-20, 2016

the formation of better crosslinked network structure than EBN cured resins.

5. References 1. Lee, G. S., Lee, Y. C., & Gong, M. S. (2001). Prepration of epoxy resins containing ether ether sulfone unit and thermal properties. Bulletin of Korean Chemical Society, 22(12), 13931396. 2. Romao, B. M. V., Diniz, M. F., Azevedo, M. F. P., Lourenco, V. L., Pardini, L. C., Dutra, R. C. L., & Burel, F. (2006). Characterization of the curing agents used in epoxy resins with TG/FT-IR technique. Polímeros: Ciencia e Technologia, 16(2), 94-98. http://dx.doi.org/10.1590/S0104-14282006000200007. 3. Paluvai, N. R., Mohanty, S., & Nayak, S. K. (2014). Synthesis and modification of epoxy resins and their composites: A review. Polymer-Plastics Technology and Engineering, 53(16), 17231758. http://dx.doi.org/10.1080/03602559.2014.919658. 4. Tao, Z., Yang, S., Ge, Z., Chen, J., & Fan, L. (2007). Synthesis and properties of novel fluorinated epoxy resins based on 1,1-bis(4-glycidyllesterphenyl)-1-(3′-trifuoromethylphenyl)2,2,2-trifluoroethane. European Polymer Journal, 43(2), 550-560. http://dx.doi.org/10.1016/j.eurpolymj.2006.10.030. 5. Vinicius, P., Bluma, S. G., & Raquel, M. S. (2013). Influence of the polyhedral oligomeric silsesquioxane n-phenylaminopropyposs in the thermal stability and the glass transition temperature of epoxy resin. Polímeros: Ciencia e Technologia, 23(3), 331-338. http://dx.doi.org/10.4322/polimeros.2013.039. 6. Alhousami, M. H. M., Al-Kamali, A. S. N., & Athawale, A. A. (2014). Synthesis and characterization of novel sulphanilamide/ epoxy resin modified polyester for thermal stability and impact strength. Open Journal of Polymer Chemistry, 4(4), 115-127. http://dx.doi.org/10.4236/ojpchem.2014.44013. 7. Dixit, V., Nagpal, A. K., & Singhal, R. (2010). Synthesis and characterization of phenoxy modified epoxy blends. Malaysian Polymer Journal, 5(2), 69-83. Retrieved in 20 December 2014, from http://www.fkkksa.utm.my/mpj 8. Oliveira, A., Backer, C. M., & Amico, S. C. (2015). Evaluation of the characteristics of an epoxy resin with diffrent degassing agents. Polímeros: Ciência e Tecnologia, 25(2), 186-191. http:// dx.doi.org/10.1590/0104-1428.1661. 9. Yu, M., Feng, B., Xie, W., Fang, L., Li, H., Liu, L., Ren, M., Sun, J., Zhang, J., & Hu, H. (2015). The modification of a terafunctional epoxy and its curing reaction. Materials, 8(6), 3671-3684. http://dx.doi.org/10.3390/ma8063671. 10. Ge, Z., Tao, Z., LiuJ., Fan, J., & Yang, S. (2007). Synthesis and charaterization of novel trifunctional fluorine containing epoxy resins based on 1,1,1-Tris(2,3-epoxypropoxyphenyl)2,2,2-trifluoroethane. Polymer Journal, 39(11), 1135-1142. http://dx.doi.org/ 10.1295/polymj.PJ2007096. 11. Cheng, J., Li, J., & Zhang, J. Y. (2009). Curing behavior and thermal properties of trifunctional epoxy resin cured by 4, 4′-diaminodiphenyl sulfone. Express Polymer Letters, 3(8), 501-509. http://dx.doi.org/10.3144/expresspolymlett.2009.62. 12. Heo, G. Y., & Park, S. J. (2009). Effect of substituted trifluoromethyl groups on thermal and mechanical properties of fluorine containing epoxy resin. Macromolecular Research, 17(11), 870-873. http://dx.doi.org/10.1007/BF03218628. 13. Meenakshi, K. S., Pradeep, E., Sudhan, J., & Kumar, S. A. (2012). Development and characterization of new phosphorus based flame retardant tetraglycidyl epoxy nanocomposites for aerospace application. Bulletin of Materials Science, 35(2), 129-136. http://dx.doi.org/10.1007/s12034-012-0271-0. 14. Liu, J. G., He, M. H., Li, Z. X., Qian, Z. G., Wang, F. S., & Yang, S. Y. (2002). Synthesis and characterization of organosoluble polyimides with trifluoro-substituted benzene in the side chain. 19


Kiran, V., & Gaur, B. Journal of Polymer Science. Part A, Polymer Chemistry, 40(10), 1572-1582. http://dx.doi.org/10.1002/pola.10240. 15. Knoll, D. W., Nelson, D. H. & Keheres, P. W. (1958). Paint, plastics and printing ink chemistry. In 134th American Chemical Society Meeting (Paper No. 5, pp. 20). Chicago: Division of Paint, Plastics and Printing Ink Chemistry. 16. Maka, H., & Spychaj, T. (2012). Epoxy resin crosslinked with conventional and deep eutectic ionic liquids. Polimery, 57(6), 456-462. http://dx.doi.org/10.14314/polimery.2012.456. 17. Gaceva, G. B., & Buzarovska, A. (2013). A rapid method for the evaluation of cure kinetics of thermosetting polymers. Macedonian Journal of Chemistry and Chemical Engineering, 32(2), 337-344. Retrieved in 20 December 2014, from http:// www. researchgate.net/260245399 18. Costa, L. M., Pardini, C. L., & Rezende, C. M. (2005). Influence of aromatic amine hardness in the cure kinetics of an epoxy resin used in advanced composites. Materials Research, 8(1), 65-70. http://dx.doi.org/10.1590/S1516-14392005000100012. 19. Su, W. F., Huang, H. W., & Pan, P. W. (2002). Thermal properties of rigid rod epoxies cured with diaminodiphenyl sulfone and dicyandiamide. Thermochimica Acta, 392/393, 391-394. http:// dx.doi.org/10.1016/S0040-6031(02)00125-9. 20. Freeman, E. S., & Carroll, B. (1958). The application of thermoanalytical techniques to reaction kinetics: the

20

thermogravimetric evaluation of the kinetics of the decomposition of calcium oxalate monohydrate. Journal of Physical Chemistry, 62(4), 394-397. http://dx.doi.org/10.1021/j150562a003. 21. Coats, A. W., & Redfern, J. (1964). Kinetic parameters from thermogravimetric data. Nature, 201(4914), 68-69. http:// dx.doi.org/10.1038/201068a0. 22. Horowitz, H. H., & Metzger, G. (1963). A new analysis of thermogravimetric traces. Analytical Chemistry, 35(10), 14641468. http://dx.doi.org/10.1021/ac60203a013. 23. Dharwadkar, S. R., Kharkhanawala, M. D., Schwenker, R. P., & Garn, P. D. (1969). Thermal analysis in organic materials and physical chemistry. New York: Academic Press. 24. Tripathi, G., & Srivastava, D. (2011). Study on the effect of carboxyl terminated butadiene acrylonitrile (CTBN) copolymer concentration on the decomposition kinetics parameters of blends of glycidyl epoxy and non-glycidyl epoxy resin. International Journal of Organic Chemistry, 1(3), 105-112. http://dx.doi.org/10.4236/ijoc.2011.13016.

Received: Dec. 20, 2014 Revised: Aug. 15, 2015 Accepted: Aug. 17, 2015

PolĂ­meros, 26(1), 11-20, 2016


http://dx.doi.org/10.1590/0104-1428.2276

Cardanol-based thermoset plastic reinforced by sponge gourd fibers (Luffa cylindrica) André Leandro da Silva1*, Lucas Renan Rocha da Silva1, Isabelle de Andrade Camargo2, Deuber Lincon da Silva Agostini3, Derval dos Santos Rosa2, Diego Lomonaco Vasconcelos de Oliveira1, Pierre Basílio Almeida Fechine4 and Selma Elaine Mazzetto1 Department of Organic and Inorganic Chemistry, Universidade Federal do Ceará – UFC, Fortaleza, CE, Brazil 2 Center for Engineering, Modeling and Applied Social Sciences, Universidade Federal do ABC – UFABC, Santo André, SP, Brazil 3 Department of Physics, Chemistry and Biology, Universidade Estadual Paulista – UNESP, Presidente Prudente, SP, Brazil 4 Department of Analytical Chemistry and Physical Chemistry, Universidade Federal do Ceará – UFC, Fortaleza, CE, Brazil 1

*andre.leandro@ufcg.edu.br

Abstract A growing global trend for maximum use of natural resources through new processes and products has enhanced studies and exploration of renewable natural materials. In this study, cardanol, a component of the cashew nut shell liquid (CNSL), was used as a building block for the development of a thermosetting matrix, which was reinforced by raw and modified sponge gourd fibers (Luffa cylindrica). DSC and TG results showed that among biocomposites, the one reinforced by sponge gourd fibers treated with NaOH 10 wt% (BF10) had the highest thermal stability, besides the best performance in the Tensile testing, showing good incorporation, dispersion, and adhesion to polymer matrix, observed by SEM. After 80 days of simulated soil experiments, it has been discovered that the presence of treated fiber allowed better biodegradability behavior to biocomposites. The biobased thermoset plastic and biocomposites showed a good potential to several applications, such as manufacturing of articles for furniture and automotive industries, especially BF10. Keywords: biobased plastic, biodegradable, composites, lignocellulosic fibers.

1. Introduction A growing global trend for maximum use of natural resources through new processes and products has enhanced researches and exploration of renewable natural materials[1]. Cashew nut shell liquid (CNSL) is a remarkable example of a renewable resource often used in polymeric materials syntheses. It is a viscous liquid in the soft honeycomb of the cashew nut, which is the crop of the cashew tree, Anacardium occidentale L, native to Brazil[2]. Anacardic acid, cardanol, cardol, and 2-methylcardol are phenolic compounds with long side chain substitution at the meta position and the main components of CNSL. These phenolic compounds can be used for producing a number of polymeric substances of wide utility in several industrial applications. The use of CNSL instead of phenol is an excellent example of conservation of a synthetically derived substance and the utilization of a byproduct[3]. Commercial grade CNSL, also known as technical CNSL, hardly contains any anacardic acid, because of decarboxylation during the roasting process, which converts anacardic acid into cardanol[4]. Cardanol is a mixture of compounds with a changeable number of unsaturations in the fifteen carbon atom chains, which is in the meta position to the hydroxyl group[5]. Cardanol-based polymers and polymeric materials are

Polímeros, 26(1), 21-29, 2016

both documented in the literature[6], as well as the focus of continuing researches. Among the CNSL constituents, cardanol is a viable option to polymeric materials synthesis, because of this compound it is a monosubstituted phenol, which decreases the problem of steric hindrance during the polymerization reaction. There is a great interest in the use of composite materials from polymer matrices reinforced by natural fibers[7]. Environment-friendly materials, their development and cleaner production processes have been needed to preserve the environment and guarantee improved living conditions for future generations. Lignocellulosic fibers are natural biopolymers constituted primarily of cellulose, hemicelluloses, and lignin, which contribute differently to the mechanical strength of the composite​​[8,9]. The application of natural fibers as a reinforcement provides composite materials with low density, higher porosity, satisfactory values of tensile strength, and impact resistance, besides greater control of fissuration and ductile behavior at break[10-12]. Natural fibers present advantages such as low cost, excellent resistance to solvents, and to temperature. In addition, they are non-toxic, non-abrasive, and easily modified by chemicals[13]. These characteristics make fibers feasible

21

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


Silva, A. L., Silva, L. R. R., Camargo, I. A., Agostini, D. L. S., Rosa, D. S., Oliveira, D. L. V., Fechine, P. B. A., & Mazzetto, S. E. materials for many applications such as in carpets, vases, cords, roof tiles, automotive upholsteries, mattresses, etc[14]. Lignocellulosic fibers have been widely investigated and used as reinforcement in polymer matrices such as coconut fibers[15], jute[16], sisal[17], banana tree[18]. Sponge gourd fiber (Luffa cylindrica) is an interesting raw material and it has not yet had its potentialities fully explored. It is a subtropical plant, abundant in China, Japan, and other Asian countries as well as countries in Central and South America. Sponge gourd has a netting-like fibrous vascular system, which is thin, resistant, viscoelastic, and soft, used mainly as a bath sponge[19]. The chemical composition of sponge gourd fibers depends on several factors such as plant origin, weather conditions, soil nature, etc. For instance, cellulose content varies from 55% to 90%, lignin content from 10% to 23%, and hemicellulose content from 8% to 22%[20,21]. However, lignocellulosic fibers are hydrophilic and because of it there is low compatibility with the polymer matrix, which is mostly hydrophobic[7]. This lack of affinity results in a low surface of adhesion, which affects the transference of effective tensions, compromising mechanical properties of composites[22,23]. Therefore, a pretreatment is necessary before using these fibers as a reinforcement in polymer matrices. Alkaline treatment is a known method to surface modification of fibers and can provide better adhesion. The main motivation of this research was due to the fact that any study has not been reported in the literature on the preparation and characterization of a thermoset plastic based on cardanol and reinforced by raw and treated sponge gourd fibers so far. The termoset plastic and the biocompostes were investigated by techniques of Gas Chromatography–Mass Spectrometry (GC-MS), Differential Scanning Calorimetry (DSC), Thermogravimetry (TG), Tensile testing, Scanning Electron Microscopy (SEM) and Biodegradation in simulated soil. In addition, this paper has also included a study on the crosslinking degree and cure conditions by a low-cost gravimetric method.

2. Materials and Methods 2.1 Sponge gourd fibers, technical CNSL and chemical reagents Sponge gourd fibers were purchased commercially (São Paulo - SP, Brazil) without any skin. Sodium hydroxide (NaOH - 97% - Dynamic) and a commercial sodium hypochlorite solution (NaClO - 2.0 wt% of active chlorine) were used for alkaline treatment and bleaching of fibers. The technical CNSL used in this study was kindly provided by Amêndoas do Brasil Ltda (Fortaleza, Brazil). Methanol (CH3OH - 99.8% - Dynamic), ammonium hydroxide (NH4OH - 28-30% - Dynamic), hexane (C6H14 - 98.5% - Dynamic), anhydrous sodium sulfate (Na2SO4 - 99% - Dynamic), hydrochloric acid (HCl - 37% - Dynamic), formaldehyde (CH2O - 36.5-38% - Dynamic), diethylenetriamine - DETA (C4H13N3 - 99% - Sigma-Aldrich), toluene (C7H8 - 99.5% - Dynamic) were used without any prior treatment. Resin epoxy DGEBA (Bisphenol A diglycidyl ether) was obtained from Avipol (São Paulo, Brazil). 22

2.2 Alkaline treatment Simple procedures were conducted prior to the sponge gourd fibers such as cutting, seeds removal, milling, and drying. Previously dried fibers were ground in a Wiley mill (FORTINOX–STAR FT 80) and fractions until 30 mesh were collected. Sponge gourd fibers were treated with sodium hydroxide solutions 5 wt% and 10 wt% at 65 °C for 4 h. Fibers were washed with distilled water to remove excess of NaOH until neutral pH was achieved. Thereafter, fibers were treated with sodium hypochlorite solution 1 wt% of active chlorine at 65 °C for 1 h. Eventually they were washed with distilled water to remove excess of NaClO and dried in an air circulation oven at 65 °C.

2.3 Prepolymer resol synthesis Prepolymer resol synthesis was carried out in a laboratory glass reactor equipped with a stirrer, thermometer and reflux condenser. Cardanol was previously isolated from technical CNSL by the method proposed by Kumar et al[24]. Resol was synthesized by mixing cardanol and formaldehyde aqueous solution, and ammonium hydroxide as a catalyst. The mixture was heated until 80 °C for 3 h 30 min under magnetic stirring. Cardanol/Formaldehyde with a molar ratio of 1:2 was prepared in this synthesis.

2.4 Termoset plastic and biocomposites preparation The thermoset plastic (Figure 1b-d) was prepared by mixing the prepolymer resol and epoxy resin as a modifier agent, and diethylenetriamine as a catalyst in the proportion 0.93:1:0.11 (wt), respectively. The use of aliphatic amines as catalysts enables materials to be produced under normal conditions of temperature and pressure, which facilitates their processing. The mixture was transferred to a metal mold with 180 mm high, 160 mm wide and, 5 mm thick (Figure 1a), then cured at ambient temperature for 48 h in a dispensing hood and post-cured in the oven at 100 °C for 4 h. Carnauba wax was used as a mold release agent. The cardanol-based resin was used as a polymer matrix and biocomposites (Figure 1e-g) were prepared by using treated and untreated sponge gourd fibers as dispersed phase (Matrix/Fiber with weight ratio of 85:15). The biocomposites processing conditions were the same used for the thermoset plastic.

2.5 Crosslinking degree determination A gravimetric technique based on ASTM D2765[25] was used to crosslinking degree investigation. Samples of the thermoset plastic were weighed (Wi) and then immersed in toluene for 24 h to extract the soluble content and determine the crosslinking degree (CD). At the end of this period, samples were dried at ambient temperature to constant weight and again weighed (Wf). The experiments were performed in triplicate for the thermoset plastic cured at ambient temperature for 48 h, as well as post-cured at 60, 80, and 100 °C for 2 h. After that, another experiment was conducted, in which the temperature was fixed and the CD was measured to post-cure ranges of 30, 60, 120, 180, 240, 300, and 360 min. The CD was calculated by the Equation 1 Polímeros, 26(1), 21-29, 2016


Cardanol-based thermoset plastic reinforced by sponge gourd fibers (Luffa cylindrica)

Figure 1. (a) Metal mold; (b) thermoset plastic after demolding; (c) thermoset plastic cross section; (d) thermoset plastic after cutting, and biocomposite after cutting reinforced by (e) untreated sponge gourd fiber, and fiber treated with NaOH; (f) 5 wt% and (g) 10 wt%.

below. The toluene used as solvent was chosen based on prior solubility tests with the thermoset plastic. Mf CD ( % = ×100 (1) ) Mi

2.6 Gas chromatography-Mass spectrometry (GS-MS) GC-MS analysis of cardanol was conducted in a QP 2010 chromatographer (Shimadzu), equipped with a DB-5 column (phenyl methylpolysiloxane 5%). The injected volume of each sample was about 1 μl and helium was used as a carrier gas. The total pressure, total flow and the split ratio were 58 KPa, 87.4 ml/min, and 100, respectively. The ion source temperature was 230 °C, the interface temperature was 300 °C, and the mass ratio of detection was 18-800.

2.7 Thermogravimetry (TG) TG curves were obtained in a TG STA 6000 equipment (Perkin Elmer) under N2 atmosphere (flow rate of 50 ml/min), heating rate of 10 °C/min in a temperature range between 30 and 600 °C to evaluate thermal stability of the thermoset plastic and biocomposites. Platinum pan was used and the weight of each sample was about 3 mg. Polímeros, 26(1), 21-29, 2016

2.8 Differential scanning calorimetry (DSC) DSC curves were obtained in a DSC Q-20 equipment (TA Instruments) under N2 atmosphere (flow rate of 50 ml/min), heating rate of 10 °C/min in a temperature range between 30 °C and 400 °C to study cure degree and look at how heat capacity of the thermoset plastic and biocomposites is changed by temperature. Hermetic aluminum pan was used and the weight of each sample was about 2.5 mg.

2.9 Tensile testing Tensile testing was performed on the thermoset plastic and biocomposites in a DL 10000/700 tensile machine (Emic) with 100 kN of load cell. The test speed was 5 mm/min and specimens were prepared with 175.0 mm high, 25 mm wide, and 5 mm thick. A minimum of five samples were tested to each material and the average value was calculated. This test was conducted according to ASTM D638[26] and ASTM D3039[27].

2.10 Scanning eletron microscopy (SEM) The thermoset plastic and biocomposites were previously coated with a thin layer of metallic gold in a sputter coater Q150R ES (Quorum). An electron microscope SSX-550 23


Silva, A. L., Silva, L. R. R., Camargo, I. A., Agostini, D. L. S., Rosa, D. S., Oliveira, D. L. V., Fechine, P. B. A., & Mazzetto, S. E. (Shimadzu) was used to get micrographs from fractured surface of specimens after Tensile testing at an accelerating voltage of 15 kV.

2.11 Biodegradation in simulated soil The biodegradation was investigated by submitting the samples to simulated soil. They were weighed and buried at 25 ± 3 °C. The simulated soil was prepared with 23% of loamy silt, 23% of organic matter (cow manure), 23% of sand, and 31% of distilled water (all wt/wt). Biodegradation was monitored for 80 days by measuring the mass variation. The buried specimens were recovered, washed with distilled water, and dried at ambient temperature until there was no further weight variation, and then weighed. After weighing, the specimens were buried again. The experiments were performed in quadruplicate.

The results of the crosslinking degree experiments were verified by DSC. It is one of the most used techniques to study the cure of thermoset materials. Figure 4 shows the DSC curve obtained by monitoring the cure reaction of the thermoset plastic. The cure reaction showed a large exothermic process at 87.76 °C and total heat of 92.95 J/g. The endothermic event

3. Results and Discussion The isolated cardanol was verified by GC-MS and its purity was confirmed (66.25% of monounsaturated, 29.04% di-unsaturated, and 4.71% saturated). Modifications in sponge gourd fibers after chemical treatment have been already discussed in earlier studies[28,29]. In all chemical treatments of this study the same concentration of NaClO (1 wt%) was used, that is why in following discussions only the concentrations of NaOH, which varied (5 and 10 wt%), will be mentioned. It is important to study the degree of crosslinking of thermoset materials, because it is linked directly to the properties of the final material. Figure 2 shows the results of crosslinking degree obtained by the gravimetric technique based on the swelling experiments to the thermoset plastic. 0 °C represents the thermoset plastic which was cured at ambient temperature for 48 h and it did not undergo any post-cure process. After this period, the material presented a crosslinking degree of 88.53% ± 0.34. Besides being cured at ambient temperature for 48 h, thermoset plastic samples were also post-cured for 2 h in temperature ranges of 60, 80, and 100 °C in which crosslinking degree was of 93.54 ± 0.84, 93.57 ± 0.36, and 95.67% ± 0.40, respectively. Temperatures over 100 °C were tested as well, but they deformed the materials. Subsequently, the temperature of post-cure was fixed at 100 °C, which showed the highest crosslinking degree to the studied time interval, and the crosslinking degree was investigated to different time ranges, as shown in Figure 3 below. The crosslinking degree to post-cure time ranges of 30, 60, 120, 180, 240, 300, and 360 min was 92.45 ± 0.45, 94.04 ± 0.95, 94.48 ± 0.80, 95.53 ± 0.26, 96.66 ± 0.14, 96.68 ± 0.22, and 96.88% ± 1.03, respectively. The crosslinking degree increased with increasing temperature and it did not show any significant variation from 240 minutes. Considering the results of the experiments, the cure and post-cure conditions to the thermoset plastic were defined to be 48 h at ambient temperature in dispensing hood and 4 h at 100 °C in oven, respectively. These conditions were also tested to biocomposites and the cure time of 48 h in dispensing hood was to guarantee total liberation of the used catalyst (DETA). 24

Figure 2. Crosslinking degree to different ranges of post-cure temperature.

Figure 3. Crosslinking degree to different ranges of post-cure time.

Figure 4. DSC curve to the cure reaction of thermoset plastic. Polímeros, 26(1), 21-29, 2016


Cardanol-based thermoset plastic reinforced by sponge gourd fibers (Luffa cylindrica) at 136 °C is because of dehydration. After investigations on the thermoset plastic, from this point forward biocomposites will also be discussed. For convenience, thermoset plastic, biocomposite reinforced by untreated sponge gourd fibers, and those reinforced by sponge gourd fibers treated with NaOH 5 and 10 wt% were labeled as TP, BFB, BF5, and BF10, respectively. The conversion degree of the TP and biocomposites were studied by DSC, which curves are shown in Figure 5. All materials presented a variation in baseline around 60 °C. This variation was attributed to glass transition temperature (Tg), which was calculated by the method based on ASTM D3418[30]. Maffezzoli et al. studied biocomposites with matrix based on cardanol and reinforced by natural fibers. They found a Tg of 56 °C to their materials[31]. There was not any exothermic event related to residual heat for the TP or any biocomposites. It indicated that cure conditions applied to the TP were also appropriate to the biocomposites, providing complete conversion, and consequently complete cure of all materials. These DSC results on materials cure validated those observed in the crosslinking degree experiments.

A short endothermic event was observed at 169, 172, and 187 °C to BFB, BF5, and BF10, respectively. It was assigned to dehydration and not observed to TP due to its hydrophobicity. The deep endothermic event at 336, 286, 294, and 315 °C refers to degradation of TP, BFB, BF5, and BF10, respectively. The TP is a little more thermally stable than biocomposites. However, chemical treatment increased the thermal stability of the biocomposites, especially BF10. It happened because the treatment partially removes lignin and hemicellulose, which start to degradate at low temperature ranges. Chemical treatment also removes surface waxes, promoting better adhesion between biocomposite phases. The better the interaction between fiber and matrix, the​​ higher is the initial temperature of thermal degradation. The weight variation of the materials as a function of temperature was studied by Thermogravimetry. TG and DTG (Differential Thermogravimetry) curves to the thermoset plastic and biocomposites are shown in Figure 6. TG curves (Figure 6a) showed a small weight loss from room temperature to 160 °C of 2.12, 2.84, and 3.84% to BFB, BF5, and BF10, respectively. This process corresponds to loss of water and it was not observed to TP because of its hydrophobicity, as also seen in DSC curves. The water content in the biocomposites increased with increased concentration of NaOH. It occurred because alkaline treatment exposed the hydroxyl groups of the fiber, which increased the interaction with water. Modibbo et al. reported alkaline treatment of fibers increases the moisture uptake[32]. DTG curves (Figure 6b) showed two stages of weight loss. In the first one, TP, BFB, BF5, and BF10 maintained thermal stability and maximum temperature peak up to about 338 and 379, 316 and 365, 324 and 366, 326 and 367 °C, respectively. This event is related to the decomposition of lignin, hemicellulose, and cellulose[33-35] and degradation of soft segments of the polymer matrix. The second stage is attributed to degradation of carbonaceous residues. At the end of these two stages, TP, BFB, BF5, and BF10 showed 14.77, 18.17, 15.48, and 14.55% of residues, respectively. TG and DTG data agree with DSC results.

Figure 5. DSC curves to TP, BFB, BF5, and BF10.

Biocomposites mechanical behavior depends on synergy between the reinforcing agent and the matrix. Table 1

Figure 6. (a) TG and (b) DTG curves to TP, BFB, BF5, and BF10. Polímeros, 26(1), 21-29, 2016

25


Silva, A. L., Silva, L. R. R., Camargo, I. A., Agostini, D. L. S., Rosa, D. S., Oliveira, D. L. V., Fechine, P. B. A., & Mazzetto, S. E. shows mechanical properties to the thermoset plastic and biocomposites. All materials were examined to study the influence of the sponge gourd fibers acting as reinforcement as well as the effects of chemical treatment on mechanical properties of biocomposites. Remarkable changes were observed. There was an increase of 97.67, 89.86, and 144.38% in tensile strength of BFB, BF5, and BF10, respectively, when compared to TP. Young modulus and breaking load values to BFB, BF5, and BF10 showed an increase of 126.58 and 118.14, 124.27 and 102.93, 138.15 and 144.12%, respectively, when compared to TP. All biocomposites presented superior mechanical properties compared to TP, especially BF10. As known, lignocellulosic fibers can be an efficient reinforcement as long as adequate is their interaction with the polymer matrix. These results indicated that sponge gourd fibers treated with NaOH 10 wt% had better interaction with polymer matrix, allowing a greater Table 1. Mechanical properties to the TP and biocomposites. Tensile strength (MPa)

Young’s modulus (GPa)

Breaking load Deformation (N) (%)

TP

7.30 ± 0.55

1.73 ± 0.23

232.74 ± 38.82

0.86

BFB

14.43 ± 1.92

3.92 ± 0.25

507.70 ± 38.04

0.77

BF5

13.86 ± 1.82

3.88 ± 0.41

472.32 ± 75.20

0.75

BF10

17.84 ± 4.08

4.12 ± 0.28

568.17 ± 136.30

1.05

stiffness, better distribution of tension and therefore, better Tensile testing results. DSC and TG data corroborated these results. After Tensile testing, micrographs were recorded from fractured surface of specimens to investigate morphology and interaction between polymer matrix and sponge gourd fibers. Figure 7 shows the micrographs to the TP and biocomposites. Micrographs showed voids on the fractured surface of TP (Figure 7a), probably because of cure reaction, which is exothermic, air incorporation or water presence. The surface of BFB (Figure 7b) and BF5 (Figure 7c) showed adhesive fracture, characterized by interface debonding between fiber and matrix due to low interaction and poor adhesion. On the other hand, BF10 surface (Figure 7d) did not present any voids or debonding, suggesting good compatibility between sponge gourd fibers and polymer matrix, which led to a better distribution of tension. Figure 8 shows the fractured region and lays emphasis on debonding in BFB as well as better adhesion and flatter fracture in BF10 at macroscopic level. It is known that pretreatment can enable better adhesion between matrix and fiber[36,37]. Micrographs and the photograph showed that the alkaline treatment provided better interface between sponge gourd fibers and cardanol-based resin, especially that one with NaOH 10 wt%. These results agree with those obtained to DSC, TG, and Tensile testing.

Figure 7. Micrographs of fractured surface to (a) TP (50×, 200 µm); (b) BFB (50×, 500 µm); (c) BF5 (50×, 500 µm) and (d) BF10 (50×, 200 µm). 26

Polímeros, 26(1), 21-29, 2016


Cardanol-based thermoset plastic reinforced by sponge gourd fibers (Luffa cylindrica)

Figure 8. Fractured region after Tensile testing to (a) BFB and (b) BF10.

There is a growing search for more durable materials during use and less damage to environment when disposed. The results of the behavior of the TP and biocomposites in simulated soil after 80 days are in Figure 9. The weight increase in TP, BFB, and BF5 in the first days occurred due to absorption or adsorption of water by these materials, which were more resistant to hydrolysis[38] when compared to BF10. TP showed the greatest water incorporation because of the voids in this material, as seen by SEM. The weight increase of the materials can also be attributed to growth of microorganisms involved in polymer degradation. During the investigated period of time (80 days), a weight loss of 10 ± 36.52, 41 ± 49.15, 65 ± 24.39 and 66% ± 13.28% to TP, BFB, BF5, and BF10, respectively, was observed. Polymers are degraded in biological systems by oxidation and hydrolysis. The first stage of aging in simulated soil corresponds to the abiotic phase, in which macromolecules hydrolyze and smaller molecules such as monomers and oligomers are formed[39]. According to Sun and Tomkinson[40], lignin protects the fiber from microorganisms action. Given this, it is possible to say that alkaline treatment was the responsible for the greatest weight decreases in BF5 and BF10 due to the removal of lignin content, which made the treated Polímeros, 26(1), 21-29, 2016

Figure 9. Biodegradation of TP and biocomposites.

fibers and their respective biocomposites more susceptible to biodegradation.

4. Conclusions The effective use of cardanol as a building block for the development of a thermosetting matrix reinforced by sponge gourd fibers was investigated. A combination of cure 27


Silva, A. L., Silva, L. R. R., Camargo, I. A., Agostini, D. L. S., Rosa, D. S., Oliveira, D. L. V., Fechine, P. B. A., & Mazzetto, S. E. at ambient temperature and post-cure were conducted to a higher handling time and better polymer matrix wettability as well as maximum crosslinking degree and complete cure. Although it does not replace classical instrumental techniques, the gravimetric method used to determine the crosslinking degree to the thermoset plastic showed that DETA can provide a satisfactory cure degree at room temperature and It was found to be an interesting alternative, and an indexing of tool for this polymer system, especially due to its low cost and easy performance. While that method supplied information on crosslinking degree, DSC results confirmed total conversion of groups involved in crosslinking and endorsed that cure and post-cure conditions applied to the thermoset plastic were appropriate to the biocomposites as well. DSC curves showed a Tg around 60 °C to the thermoset plastic and biocomposites, which means sponge gourd fibers did not affect the Tg. Although the thermoset plastic has had a little higher thermal stability than biocomposites, sponge gourd fibers seemed to be a viable reinforcement, especially due to extraordinary improvement in mechanical properties of the biocomposites. DSC and TG results showed that among biocomposites, that one reinforced by sponge gourd fibers treated with NaOH 10 wt% (BF10) had the highest thermal stability. BF10 also presented the best performance in Tensile testing and good incorporation, dispersion, and adhesion with polymer matrix, observed by SEM. After 80 days of simulated soil experiments was found that the presence of treated fiber allowed better biodegradability behavior to biocomposites. This paper provided useful information and introduced the experimental set of results on the synthesis and characterization of a thermoset plastic and biocomposites prepared from a combination between a cardanol-based resin and sponge gourd fibers by a low-cost processing, which showed a good potential to several applications, from manufacturing of articles to furniture and automotive industries.

5. References 1. John, M. J., & Thomas, S. (2008). Biofibres and biocomposites. Carbohydrate Polymers, 71(3), 343-364. http://dx.doi. org/10.1016/j.carbpol.2007.05.040. 2. Harvey, M. T., & Caplan, C. (1940). Cashew nut shell liquid. Industrial & Engineering Chemistry, 3(10), 1306-1310. http:// dx.doi.org/10.1021/ie50370a008. 3. Manjula, S., Sudha, J. D., Bera, S. C., & Pillai, C. K. S. (1985). Polymeric resin from renewable resources: studies on polymerization of the phenolic component of coconut shell tar. Journal of Applied Polymer Science, 30(4), 1767-1771. http://dx.doi.org/10.1002/app.1985.070300440. 4. Tyman, H. P., Wilczynski, D., & Kashani, M. A. (1978). Compositional studies on technical cashew nutshell liquid (CNSL) by chromatography and mass spectroscopy. Journal of the American Oil Chemists’ Society, 55(9), 663-668. http:// dx.doi.org/10.1007/BF02682455. 5. Mohapatra, S., & Nando, G. B. (2014). Cardanol: a green substitute for aromatic oil as a plasticizer in natural rubber. Royal Society of Chemistry Advances, 4, 15406-15418. http:// dx.doi.org/10.1039/c3ra46061d. 6. Balachandran, V. S., Jadhav, S. R., Vemula, P. K., & John, G. (2013). Recent advances in cardanol chemistry in a nutshell: from a nut to nanomaterials. Chemical Society Reviews, 28

42(2), 427-438. http://dx.doi.org/10.1039/C2CS35344J. PMid:23114456. 7. Corrales, F., Vilaseca, F., Llop, M., Gironès, J., Méndez, J. A., & Mutjè, P. (2007). Chemical modification of jute fibers for production of green-composites. Journal of Hazardous Materials, 144(3), 730-735. http://dx.doi.org/10.1016/j. jhazmat.2007.01.103. PMid:17320283. 8. Zou, L., Jin, H., Lu, W., & Li, X. (2009). Nanoscale structural and mechanical characterization of the cell wall of bamboo fibers. Materials Science and Engineering C, 29(4), 1375-1379. http://dx.doi.org/10.1016/j.msec.2008.11.007. 9. Rao, K. M. M., & Rao, K. M. (2007). Extraction and tensile properties of natural fibers: vakka, date and bamboo. Composite Structures, 77(3), 288-295. http://dx.doi.org/10.1016/j. compstruct.2005.07.023. 10. Kumar, V., Kushwaha, P. K., & Kumar, R. (2011). Impedancespectroscopy analysis of oriented and mercerized bamboo fiber-reinforced epoxy composite. Journal of Materials Science, 46(10), 3445-3451. http://dx.doi.org/10.1007/s10853-0115249-6. 11. Sen, T., & Reddy, H. N. J. (2011). Applications of sisal, bamboo, coir and jute and natural composites in structural up gradation. International Journal of Inovation Management and Technology, 2(3), 186-191. http://dx.doi.org/10.7763/IJIMT.2011.V2.129. 12. Chandramohan, D., & Marimuthu, K. (2011). A review on natural fibers. International Journal of Research and Reviews in Applied Sciences, 8(2), 194-206. Retrieved 20 January 2015, from http://www.arpapress.com/Volumes/Vol8Issue2/ IJRRAS_8_2_09.pdf 13. Dahlke, B., Larbig, H., Scherzer, H. D., & Poltrock, R. J. (1998). Natural fiber reinforced foams based on renewable resources for automotive interior applications. Journal of Cellular Plastics, 34(4), 361-379. http://dx.doi.org/10.1177/0021955X9803400406. 14. Colom, X., Carrasco, F., Pages, P., & Canavate, J. (2003). Effects of different treatments on the interface of HDPE/lignocellulosic fiber composites. Composites Science and Technology, 63(2), 161-169. http://dx.doi.org/10.1016/S0266-3538(02)00248-8. 15. Esmeraldo, M. A., Barreto, A. C. H., Freitas, J. E., Fechine, P. B. A., Sombra, A. B. S., Corradini, E., Mele, G., Maffezzoli, A., & Mazzetto, S. E. (2010). Dwarf-green coconut fibers: a versatile natural renewable raw bioresource. Treatment, morphology, and physicochemical properties. BioResources, 5(4), 2478-2501. Retrieved 20 January 2015, from http://ojs.cnr. ncsu.edu/index.php/BioRes/article/view/BioRes_05_4_2478_ Esmeraldo_GFFSCMMM_Draft_Green_Coconut_Fibers 16. Barreto, A. C. H., Esmeraldo, M. A., Rosa, D. S., Fechine, P. B. A., & Mazzetto, S. E. (2010). Cardanol biocomposites reinforced with juta fiber: microstructure, biodegradability, and mechanical properties. Polymer Composites, 31(11), 1928-1937. http://dx.doi.org/10.1002/pc.20990. 17. Barreto, A. C. H., Rosa, D. S., Fechine, P. B. A., & Mazzetto, S. E. (2011). Properties of sisal fibers treated by alkali solution and their application into cardanol-based biocomposites. Composites Part A: Applied Science and Manufacturing, 42(5), 492-500. http://dx.doi.org/10.1016/j.compositesa.2011.01.008. 18. Barreto, A. C. H., Costa, A. E., Jr., Freitas, J. E. B., Rosa, D. S., Barcellos, W. M., Freire, F. N. A., Fechine, P. B. A., & Mazzetto, S. E. (2013). Biocomposites from dwarf-green Brazilian coconut impregnated with cashew nut shell liquid resin. Journal of Composite Materials, 47(4), 459-466. http:// dx.doi.org/10.1177/0021998312441041. 19. Mazali, I., & Alves, O. L. (2005). Morphosynthesis: high fidelity inorganic replica of the fibrous network of loofa sponge (Luffa cylindrica). Academia Brasileira de Ciências, 77(1), 25-31. 20. Satyanarayana, K. G., Guimarães, J. L., & Wypych, F. (2007). Studies on lignocellulosic fibers of Brazil. Part I: Source, Polímeros, 26(1), 21-29, 2016


Cardanol-based thermoset plastic reinforced by sponge gourd fibers (Luffa cylindrica) production, morphology, properties and applications. Composites. Part A, Applied Science and Manufacturing, 38(7), 1694-1709. http://dx.doi.org/10.1016/j.compositesa.2007.02.006. 21. Tanobe, V. O. A., Sydenstricker, T. H. D., Munaro, M., & Amico, S. C. (2007). A comprehensive characterization of chemically treated sponge-gourds (Luffa cylindrical). Polymer Testing, 24(4), 474-482. http://dx.doi.org/10.1016/j. polymertesting.2004.12.004. 22. Habibi, Y., El-Zawawy, W. K., Ibrahim, M. M., & Dufresne, A. (2008). Processing and characterization of reinforced polyethylene composites made with lignocellulosic fibers from Egyptian agro-industrial residues. Composites Science and Technology, 68(7-8), 1877-1885. http://dx.doi.org/10.1016/j. compscitech.2008.01.008. 23. Le Troedec, M. L., Sedan, D., Peyratout, C., Bonnet, J. P., Smith, A., Guinebretiere, R., Gloaguen, V., & Krausz, P. (2008). Influence of various chemical treatments on the composition and structure of hemp fiber. Composites Part A: Applied Science and Manufacturing, 39(3), 514-522. http:// dx.doi.org/10.1016/j.compositesa.2007.12.001. 24. Kumar, P. P., Paramashivappa, P. J., Vithayathil, P. J., Subra Rao, P. V., & Srinivasa, R. A. (2002). Process for isolation of cardanol from technical cashew (Anacardium occidentale.) nut shell liquid. Journal of Agricultural and Food Chemistry, 50(16), 4705-4708. http://dx.doi.org/10.1021/jf020224w. PMid:12137500. 25. American Society for Testing and Materials – ASTM. (2006). ASTM D2765-11: standard test methods for determination of gel content and swell ratio of crosslinked ethylene plastics. West Conshohocken: ASTM International. 26. American Society for Testing and Materials – ASTM. (2014). ASTM D638-14: standard test method for tensile properties of plastics. West Conshohocken: ASTM International. 27. American Society for Testing and Materials – ASTM. (2014). ASTM D3039/D3039M-14: standard test method for tensile properties of polymer matrix composite materials. West Conshohocken: ASTM International. 28. Silva, A. L., Costa, A. E., Jr., Nascimento, D. M., Rosa, M. F., Fechine, P. B. A., & Mazzetto, S. E. (2013). Efeito do tratamento alcalino e branqueamento na morfologia e no índice de cristalinidade da fibra de bucha vegetal (Luffa cylindrical). In Anais do 53º Congresso Brasileiro de Química (pp. 2). Rio de Janeiro: Associação Brasileira de Química. Retrieved 17 September 2014, from http://www.abq.org.br/cbq/2013/ trabalhos/12/2405-16576.html 29. Silva, A. L., Costa, A. E., Jr., Nascimento, D. M., Silva, M. A. S., Sombra, A. S. B., Rosa, M. F., Fechine, P. B. A., & Mazzetto, S. E. (2013). Modificações espectroscópicas vibracionais e nas propriedades dielétricas em fibras de bucha vegetal (Luffa cylindrica) após tratamento químico. In Anais do 53º Congresso Brasileiro de Química (pp. 3). Rio de Janeiro: Associação Brasileira de Química. Retrieved 17 September 2014, from http://www.abq.org.br/cbq/2013/trabalhos/12/2408-16576. html 30. American Society for Testing and Materials – ASTM. (2015). ASTM D3418-15: standard test method for transition temperatures

Polímeros, 26(1), 21-29, 2016

and enthalpies of fusion and crystallization of polymers by differential scanning calorimetry, West Conshohocken: ASTM International. 31. Maffezzoli, A., Calò, E., Zurlo, S., Mele, G., Tarzia, A., & Stifani, C. (2004). Cardanol Based Matrix Biocomposites Reinforced With Natural Fibers. Composites Science and Technology, 64(6), 839-845. http://dx.doi.org/10.1016/j. compscitech.2003.09.010. 32. Modibbo, U. U., Alyiu, B. A., Nkafamiya, I. I., & Manji, A. J. (2007). The effect of moisture imbibition on cellulosic bast fibres as industrial raw materials. Internacional Journal of Physical Science, 2(7), 163-168. Retrieved 20 October 2014, from http://www.academicjournals.org/journal/IJPS/articleabstract/69B629713202 33. Khan, A. F., & Ahmad, S. R. (1996). Chemical Modification and spectroscopic analysis of Jute fibre. Polymer Degradation & Stability, 52(3), 335-340. http://dx.doi.org/10.1016/01413910(95)00240-5. 34. Antich, P., Vázquez, A., Mondragon, I., & Bernal, C. (2006). Mechanical behavior of high impact polystyrene reinforced with short sisal fibers. Composites Part A: Applied Science and Manufacturing, 37(1), 139-150. http://dx.doi.org/10.1016/j. compositesa.2004.12.002. 35. Szczesniak, L., Rachocki, A., & Tritt-Goc, J. (2008). Glass transition temperature and thermal decomposition of cellulose powder. Cellulose (London, England), 15(3), 445-451. http:// dx.doi.org/10.1007/s10570-007-9192-2. 36. Vázquez, G., González, S., Freire, S., & Antorrena, G. (1997). Effect of chemical modification of lignin on the gluebond performance of lignin-phenolic resin. Bioresource Technology, 60(3), 191-198. http://dx.doi.org/10.1016/S0960-8524(97)000308. 37. Kharade, A. Y., & Kale, D. D. (1998). Effect of lignin on phenolic novolak resins and moulding powder. European Polymer Journal, 34(2), 201-205. http://dx.doi.org/10.1016/ S0014-3057(97)00118-3. 38. Rosa, D. S., Bardi, M. A. G., Guedes, C. G. F., & Angelis, D. A. (2009). Role of polyethylene-graft-glycidyl methacrylate compatibilizer on the biodegradation of poly (ε-caprolactone)/ cellulose acetate blends. Polymers for Advanced Technologies, 20(12), 863-870. http://dx.doi.org/10.1002/pat.1302. 39. Kyrikou, J., & Briassoulis, D. (2007). Biodegradation of agricultural plastic films: a critical review. Journal of Polymers and the Environment, 15(12), 125-150. http://dx.doi.org/10.1007/ s10924-007-0053-8. 40. Sun, R. C., & Tomkinson, J. (2002). Comparative study of lignins isolated by alkali and ultrasound-assisted alkali extractions from wheat straw. Ultrasonics Sonochemistry, 9(2), 85-93. http://dx.doi.org/10.1016/S1350-4177(01)00106-7. PMid:11794023. Received: June 20, 2015 Accepted: Aug. 21, 2015

29


http://dx.doi.org/10.1590/0104-1428.2106

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

Effect of the hardener to epoxy monomer ratio on the water absorption behavior of the DGEBA/TETA epoxy system Ayrton Alef Castanheira Pereira1* and José Roberto Moraes d’Almeida1,2 1

Mechanical Engineering Department, Universidade do Estado do Rio de Janeiro – UERJ, Rio de Janeiro, RJ, Brazil 2 Chemical and Materials Engineering Department, Pontifícia Universidade Católica do Rio de Janeiro – PUC-RJ, Rio de Janeiro, RJ, Brazil *pereira.ayrton@gmail.com

Abstract The water absorption behavior of the DGEBA/TETA epoxy system was evaluated as a function of the epoxy monomer to amine hardener ratio. Weight gain versus immersion time curves were obtained and the experimental points were fitted using Fickian and Non-Fickian diffusion models. The results obtained showed that for all epoxy monomer to hardener ratios analyzed water diffusion followed non-Fickian behavior. It was possible to correlate the water absorption behavior to the macromolecular structure developed when the epoxy/ hardener ratio was varied. All epoxy/hardener ratios present a two-phase macromolecular structure, composed of regions with high crosslink density and regions with lower crosslinking. Epoxy rich systems have a more open macromolecular structure with a lower fraction of the dense phase than the amine rich systems, which present a more compact two-phase structure. Keywords: water absorption, epoxy resins, hardener to epoxy monomer ratio.

1. Introduction Epoxy resins show high reactivity with many different chemical compounds, like aliphatic and aromatic amines, anhydrides and polyamides[1]. This characteristic is due to the presence of the very strained ethoxyline ring structure[1,2]. Therefore, given a specific epoxy monomer, the mechanical properties of the epoxy system can be varied over fairly high bounds by changing the curing agent. Therefore, the observed variations on properties reflect differences on the macromolecular network developed. Besides the chemical nature of the hardener, other variables like the time and temperature of cure[3,4] and the hardener to epoxy monomer ratio[5] could produce very different macromolecular structures. For the epoxy system formed by the difunctional epoxy monomer, diglycidyl ether of bisphenol-A (DGEBA), and the hexafunctional aliphatic amine, triethylenetetramine (TETA), the hardener to epoxy ratio was shown to strongly affect mechanical[6,7] and thermal properties[8]. However, non-stoichiometric hardener/epoxy ratios could produce unstable macromolecular networks that could age more readily, due to temperature changes or even moisture absorption. The modifications induced by both temperature and/or moisture are directly linked to the presence of unreacted sites, maintained latent during the gelation and setting of the resin. Moisture absorption, in particular, could also promote undesirable dimensional changes on finished parts. It has to be noted that even for a completely cured epoxy system, many hydrophilic sites could be present on the final network developed. In fact, for this epoxy system the cure reactions scenarios are leaded by the primary amino addition reaction, occurring between

30

primary amines (–NH2) and the epoxy group, resulting on hydrophilic hydroxyl groups (–OH)[1]. For non-stoichiometric formulations with excess of epoxy monomer the epoxy rings could, in principle react with these hydroxyls groups forming ether groups[1,4]. But this secondary reaction will not contribute to reduce the number of hydrophilic OH groups and its effect is very restricted for reactions taking place below 150 °C[9]. Besides, amines are also strongly hydrophilic groups, and when in excess could contribute to moisture up-take and to resin plasticization[10]. Earlier studies on the DGEBA/TETA system showed that the degree of cure, and therefore the presence of latent unreacted sites, is greatly affected by the hardener to epoxy ratio and that, for room temperature cured resins without post-curing, the consume of epoxy rings go to completion only for off-stoichiometric hardener rich mixtures[7]. Also, thermo-gravimetric analysis showed that some off-stoichiometric mixtures could have a two phase like microstructure, where very crosslinked domains could be embedded on a less crosslinked matrix[8]. These structural characteristics can strongly affect the moisture up-take behavior and contribute to a faster decrease of the mechanical performance of the material. Therefore, in this work, the water absorption behavior of the DGEBA/TETA epoxy system was studied as a function of the hardener to epoxy ratio. Besides the stoichiometric formulation, epoxy rich as well as hardener rich mixtures were analyzed. The results obtained were correlated with former proposed networks developed due to the change on the hardener/epoxy ratio[8], and with the presence of unreacted sites.

Polímeros, 26(1), 30-37, 2016


Effect of the hardener to epoxy monomer ratio on the water absorption behavior of the DGEBA/TETA epoxy system v2.

Theoretical Background

2.1 Fickian behavior Several models are used to describe diffusion on polymers. The most common one uses the theoretical background of the 2nd Fick equation, namely[11] ∂c ∂  ∂c  = D  ∂t ∂x  ∂x 

(1)

In Equation 1 c is the water concentration at time t, D the diffusion coefficient and x the space coordinate measured normal to the cross section. For infinity large plates of thickness h and considering that diffusion occurs only perpendicularly to the specimens’ thickness, the solution of Equation 1 is given by[11]: M% 8 ∞ = 1− ∑ M∞ π² j = 0

 2  Dt   exp  − ( 2 j + 1) π  2    h   ( 2 j + 1) ²

(2)

where M% is the mass of water absorbed at a time t, and M∞ is the mass absorbed at saturation. Equation 2 converges fast as t increases, and, therefore, one can use the first term of the series as a good approximation. An analytical simplification of Equation 2 for values of M%/M∞> 0.5 is given by[12]: 0.75  M% D t  = 1 − exp  −7.3  x2   M∞  S   

(3)

where S = h if both sides of the test specimen are exposed to the absorption medium and S = 2h if only one side is exposed. For short times of exposure, when M%/M∞< 0.5 Equation 2 can be approximated by[12]: M % 4 Dt = M∞ h π

(4)

A general equation based on the same approach of the Fick law but covering the entire range of absorption was proposed by McKague et al.[13], namely: M% M∞

 4 Dt  = tanh   h π 

(5)

regions with light crosslinking)[21,22]. Some of these models also include a term considering that stress relaxation can occur due to swelling after water uptake[23,24]. 2.2.1 Jacobs-Jones model This model hypothesizes that certain polymers present a two-phase structure, where regions with different crosslinking densities coexist, namely: a heavily crosslinked structure, here named as the dense phase, and a lightly crosslinked structure. For such materials the absorption curve shows two main regions as depicted at Figure 1. The first stage of the water absorption curve (region I) is characterized by a fast water uptake. This region is associated to diffusion of water at both phases – i.e., at the dense and at the less dense macromolecular structure. At region II, diffusion occurs more slowly, and the water uptake is attributed only to diffusion at the dense phase, since saturation has already occurred at the less crosslinked phase[21]. To model the behavior of this two-phase polymer, the equation developed by Jacobs and Jones assumed that diffusion is governed by a Fickian behavior at both phases[21,22]. The resultant equation is as follow: 0.75   M%  D t    = Vd 1 − exp  −7.3  d2    + M∞   h     0.75     Dt (1 − Vd ) 1 − exp  −7.3  2l     h     

(6)

where Dd and Dl are the diffusion coefficients at the dense and at the less dense phase, respectively. Vd is the volume fraction associated to the dense phase, and its value depends on both Dd and Dl values. The nominal diffusion coefficient, Dx , can be calculated from the initial slope of the absorption curve (mx) at region I using the following equation[21]: 2

 m h  Dx = π  x   4M ∞ 

(7)

2.2 Non-fickian behavior Several authors consider that the Fickian model is not able to describe the complete water absorption behavior of several polymers and/or polymer composites[14-17]. In fact, in many instances it was found that the Fickian behavior can fit the experimental data points only at the early stages of the absorption process, failing, however, to describe the behavior when the absorption time increases. Other models have appeared to overcome the inadequacy of the Fickian model to describe the experimental behavior of certain polymers and polymer composites models. These models take new assumptions such as water molecules acting both as a bound and unbound phase[18-20]. or considering that polymers can have phases with different macromolecular structures (for example, regions heavily crosslinked and Polímeros, 26(1), 30-37, 2016

Figure 1. Schematic water absorption curve for resins showing a two phase structure. 31


Pereira, A. A. C., & d’ Almeida, J. R. M. 2.2.2 Modified Jacobs-Jones model

2.2.4 Carter-Kibler model

The modified Jacobs-Jones model considers that the polymer structure consists of a phase where the major amount of water is absorbed (phase 1) and another phase with a different density and/or hydrophilic character (phase 2). The equation of this model is as follow: [25]

0.75    D t    M % = M1 1 − exp  −7.3  12    +   h     0.75    D t    M 2 1 − exp  −7.3  22      h    

(8)

where D1 and D2 represent the diffusion coefficients of phase 1 and of phase 2, respectively, and M1 and M2 are the saturation values of water absorbed at each phase. The maximum absorption value at saturation, M∞, equals M1 + M2. Following the same theoretical approach of the Jacobs‑Jones model where D2 is the diffusion coefficient of the dense phase, here D2 refers to the diffusion coefficient of the polar, hydrophilic phase whereas D1 is associated to the less dense phase or to the non-polar phase at the modified model. In both models D2 will be an order of magnitude smaller than D1, since water diffusion will be hindered by the high crosslinking density or by attraction of water molecules by hydrophilic polar groups[25]. 2.2.3 Berens-Hopfenberg model The Berens and Hopfenberg model[23] includes relaxation effects due to swelling into the diffusion process. This model was successfully used to describe the behavior of polymers as well as composites[16]. Its basic equation considers that the amount of absorbed water (M%) can be represented by adding a term related to Fickian diffusion process (M%, F) and a term related to relaxation (M%, R), namely: = M % M %, F + M %, R

(9)

The term representing the Fickian diffusion behavior equals Equation 3, namely: 0.75    Dt    M t ,= M ∞, F 1 − exp  −7.3  2    F    h    

(10)

where here M∞, F is the water saturation level disregarding any stress relaxation. The term related to the relaxation effect is given by:

(

(

M t ,= M ∞, R 1 − exp −kt 2 R

))

(11)

where k is a constant related to the relaxation rate of the material and M∞, R is the water saturation value related to the relaxation event. This term does not depend on the size of the diffusing molecules, and is only related to stress relief due to swelling caused by water absorption[23]. Therefore, water absorption at a time t can be written as: 0.75    Dt    M= M ∞, F 1 − exp  −7.3  2    + %    h    (12) 

(

(

M ∞, R 1 − exp −kt 2 32

))

Trying to solve problems associated to materials failing to follow the usual Fickian models, Carter and Kibler[18] proposed a model based on two hypothesis, namely: i) the diffusion coefficient does not depend on the water concentration inside the material, and ii) the water molecules coming from the diffusion process itself and/or present at the material can be divided into two phases. This model adds two new parameters, a and b These new parameters are related to the probability by unit of time that a water molecule at the free phase transforms into a bound molecule (α) or the probability by unit of time that a bound molecule becomes a free one (β)[26]. Taking into account the following boundary conditions[19]: a

Dπ 2 h

2

and b 

Dπ 2 h2

(13)

It can be written[20]:  b a 8  Dt   M= M ∞ 1 − exp [ at ] − . 2 exp  − 2   (14) % a + b a + b π  h  

3. Materials and Methods The samples were prepared by mixing proper quantities of DGEBA epoxy monomer and TETA hardener, which were weighed within ± 0.002 g. These chemicals, from Dow Chemical (Brazil), were used as received, without any further purification. Eight different hardener/resin ratios were prepared, covering the range of epoxy rich to amine rich compositions, and including the stoichiometric one. The different hardener/resin ratios used in this work were labeled according to the amount of hardener per hundred parts of resin, in weight, denoted henceforth as phr. The ratios used are referred to in Tables 1-3, and phr 13 corresponds to the stoichiometric composition. The samples were cast in plate-shaped open silicone molds, with dimensions appropriate to water absorption measurements (n = 25 mm, l = 105 mm and h = 4 mm), and were cured at room temperature, 25 ± 3 °C. The samples obtained were dried to constant weight at 60 °C and were, then, soaked in distilled water. Care was taken in order to avoid contact of the specimens with the walls of the containers. Therefore, all surfaces of the specimens were in close contact with the soaking medium. The weight gain vs. time of immersion curve was obtained following the procedures described by the ASTM D570 standard for plastics. The experimental data obtained were modeled using the theoretical models described at Section 2.

4. Experimental Results and Discussion The results obtained by fitting the models described in item 2 to the experimental points are listed in Tables 1-3. Table 1 shows the fitted values using the Fickian models described in item 2.1. Tables 2, 3 are devoted to list the adjusted values when a non-Fickian behavior was used, item 2.2. The fitted curves were obtained using non-linear regression, and the least squares method to obtain the best Polímeros, 26(1), 30-37, 2016


Effect of the hardener to epoxy monomer ratio on the water absorption behavior of the DGEBA/TETA epoxy system Table 1. Fickian models – Equations 3 and 5. PHR

M∞ (%)

7 9 11 13 15 17 19 21

2.78 3.41 3.39 4.01 5.11 5.26 5.83 6.56

Fick Model D x 10–7 (mm2/s) 3.94 2.46 3.52 2.61 1.67 2.02 2.01 1.70

r2

M∞ (%)

0.90 0.93 0.92 0.93 0.92 0.91 0.93 0.94

3.08 4.16 4.33 4.81 7.18 6.82 7.41 9.01

McKague Model D x 10–7 (mm2/s) 4.07 1.92 3.46 2.29 0.94 1.36 1.43 1.01

r2 0.95 0.97 0.96 0.97 0.97 0.96 0.97 0.98

Table 2. Non-Fickian models – Equations 6 and 8. PHR

M∞ (%)

Vd

5.05 4.97 6.37 4.82 7.74 12.04 8.13 8.91

0.8 0.8 0.8 0.9 0.9 0.9 0.9 0.9

7 9 11 13 15 17 19 21

Jacobs-Jones Model D × 10–7 Dd × 10–8 Dl × 10–5 (mm2/s) (mm2/s) (mm2/s) 2.39 3.03 0.50 1.28 4.54 0.27 0.90 2.65 0.26 2.17 12.24 18.08 0.98 4.23 6.09 0.49 1.82 0.59 0.81 6.05 13.23 0.83 5.66 10.01

r

2

0.98 0.98 0.97 0.98 0.99 0.98 0.99 0.99

M1 (%) 0.98 1.00 1.45 0.69 0.77 1.40 0.86 0.86

Modified Jacobs-Jones Model M2 M∞ D1 × 10–7 D2 × 10–9 (%) (%) (mm2/s) (mm2/s) 3.57 4.55 0.53 39.60 10.49 11.49 0.32 9.20 16.03 17.48 0.21 3.90 4.40 5.10 7.57 92.00 7.82 8.60 6.16 34.29 124.10 125.50 0.41 0.51 7.02 7.88 11.36 63.61 8.35 9.21 11.10 53.31

r2 0.98 0.98 0.97 0.98 0.98 0.98 0.99 0.99

Table 3. Non-Fickian models – Equation 12. PHR

M∞,F (%)

7 9 11 13 15 17 19 21

2.47 3.01 3.06 3.81 4.49 4.79 5.27 5.56

Berens and Hopfenberg Model M∞,R (%) M∞* (%) k × 10–12 (s–2) 0.60 0.53 0.74 0.90 0.87 0.92 1.00 1.25

3.07 3.54 3.80 4.71 5.37 5.71 6.27 6.81

5.69 1.06 1.18 0.29 0.11 0.24 0.20 0.12

D × 10–7 (mm2/s) 1.64 1.52 1.49 1.13 1.11 1.09 1.15 1.10

r2 0.96 0.95 0.96 0.96 0.95 0.94 0.96 0.96

*M∞ = M∞,F + M∞,R[24].

fit between the experimental data points and the theoretical equations. All models except the one from Carter and Kibler[18] could be fitted to the experimental results. The lack of consistence of the Carter and Kibler model can be explained regarding that the boundary conditions stablished by Equation 13 were not satisfied – i.e., the values obtained to both a and b are greater than

h2

2

. Therefore, the results obtained when this

model was applied are not included at the present topic. The values of Vd listed at Table 2 (Jacob-Jones model) were obtained taking into account the best correlation coefficient between the experimental points and the fitted theoretical curve when Vd was varied between its boundary values – from 0 to 1 – at steps of 0.1. This procedure was used because Vd depends on the values of both Dd and Dl, and, therefore, its value is a necessary condition to apply Equation 6. Polímeros, 26(1), 30-37, 2016

To verify the results listed at Table 2 for Vd, the graphical methodology described at the work of Jacobs and Jones[21,22] was also used. The complete description of the graphical approach can be found at the works of Jacobs and Jones[21,22], and is not reproduced here for the sake of shortness. As presented on several papers water absorption and diffusion on polymers are related to different factors, but are mainly affected by the free volume existent at the macromolecular structure and by the affinity of the specific polymer to water[17,27]. The amount of free volume is considered as the main driving force to water absorption, and is related to several different physical characteristics of the polymer[15,16,28-31]. These physical characteristics are intimately linked to the degree of cure, the stoichiometric ratio and with the stiffness of the molecular bonds[29-31]. The chemical affinity of a polymer to water, by the other side, is attributed to the 33


Pereira, A. A. C., & d’ Almeida, J. R. M. polymer polarity – i.e., to the presence of sites with hydrogen bonds along the polymeric chain[26]. From the experimental results, in general, it is observed that higher values of M∞ were obtained with the increase of the hardener ratio. In fact, in a previous work Soles and Yee[31] found that the increase of the amine ratio resulted on an increase of the amount of absorbed water. The increase of the water up-take was attributed to the increase of the free volume of the epoxy system, what eases the water diffusion path and increases the number of sites to be occupied by water molecules[29]. Carfagna et al.[32], using differential scanning calorimetry, has observed also that if a large excess of hardener is used high levels of water uptake are observed, and this behavior can be linked to the formation of microcavities within the matrix in an exothermic process.

Based on this fitting, it can be said that for all tested hardener/epoxy ratios a macromolecular structure with two phases was formed. Namely, a dense phase with a large number of crosslinks and a less dense phase with fewer crosslinks, as predicted by the models of Jacobs and Jones[21,22]. For all formulations tested, the proportion of the dense phase present in the system was superior to that of the “less” dense phase, what is depicted by the high volume fraction obtained for the dense phase (Vd).

While the amount of water increased with the increase of the amount of hardener, diffusivity (D) showed little variation or decreased with the increase of the hardener content. Since the values of M∞ are primarely governed by an increase on the free volume of the polymer, the decrease of the diffusivity can be explained by the increase on the polarity of the system and by topological changes on the chains[16,29]. The increase in free volume promotes growth in the number of nanovoids throughout the network, which in turn act as routes of access of water molecules to the interaction sites. This will slow molecular motion due to water affinity to hydroxyl (-OH) groups present along the chain, slowing the diffusion process[28,29].

The variation found in the proportion of the dense phase is closely related to the topology of the chains along the material. The different ways the crosslinking reactions can occur due to the variation of the hardener to epoxy monomer ratio can lead to the formation of a more open structure, contributing to the diffusion of small molecules such as water, or a more compact structure, acting as barriers to the movement of molecules[29]. For the epoxy rich systems (phr 7, 9 and 11) the formation of a more open structure is likely, because after depletion of amine groups, secondary reactions such as homopolimerization and ether formation can occur[4,9]. This structure will favor higher diffusion rates (D) and will have smaller proportion of dense phase (Vd). Amine rich systems, by contrary, have a more closed structure, since excess of amine hardener in these systems will promote opening of all available epoxy rings, and will result on a highly crosslinked structure, leading to a decrease in diffusivity (D) and an increase in the volume fraction of the dense phase (Vd).

Comparing the results obtained between Fickian and non‑Fickian models, Tables 1-3, it is verified that the latter have a best fit to the experimental data for all tested proportions of hardener. In fact, the determination coefficient values​​ (r2) for non-Fickian models were very close to 1, denoting very high correlation between the curve stipulated by the models and the experimental data. It is worth saying that the McKague model[13] also presented high correlation with experimental results. However, it is only a mathematical modification of the Fick’s approach, without including any new physical approach to the diffusion problem.

The curves obtained by applying the Jacobs-Jones models, Table 2, showed high convergence for all tested hardener to epoxy ratios. Examples of the curve fitting to the experimental points are shown in Figure 2. In some cases, however, the values for M∞ diverged and showed inconsistent results. This behavior was particularly observed for the phr 17 ratio at the Jacobs-Jones model and for phr 9, 11 and 17 ratios at the modified Jacobs-Jones model (Figure 3). The results obtained were associated to the fact that the experimental curves showed a steady increase of the amount of water, without a clear plateau indicating

Figure 2. Fitting of the experimental points to the Jacobs-Jones models. Characteristic curves for (a) phr 7, 9 and 11 (b) phr 13, 15, 17, 19 and 21. 34

Polímeros, 26(1), 30-37, 2016


Effect of the hardener to epoxy monomer ratio on the water absorption behavior of the DGEBA/TETA epoxy system

Figure 3. Characteristic curve exhibited by phr 17 for Jacobs‑Model, and phr 9, 11and 17for modified Jacobs-Jones model showing an abnormal behavior.

The good fit of the Berens-Hopfenberg model to the experimental points does not exclude the approach and discussion thus far made based on the results found by the Jacobs-Jones model, but rather complements and confirms the behavior of both D and M∞, using, however, the concept of stress relaxation. Clearly, it can be seen that the characteristic curve obtained for the different hardener ratios (Figure 4) can be divided into two regions: an initial absorption, following an almost Fickian behavior and a subsequent absorption with slower rates of weight gain. The second region reflects the changes that occur in the matrix resin as result of the stress relaxation[30]. Water molecules are generally linked to hydroxyl groups (OH) formed during the process of opening of the epoxy ring, contributing to swelling of the material and consequently to the relaxation process. Water molecules are divided into two phases: one phase bound to the polymer chain and a free phase occupying the empty spaces present within the structure of the material[17]. The higher the proportion of free phase in relation to the bonded phase, the greater the mass gain at equilibrium, since a large number of water molecules will be “loose” to fill an increased amount of free spaces. Based on the results obtained by applying the Berens‑Hopfenberg model, it can be seen that the greater the amount of amine in the epoxy system, the greater the proportion of free phase is, since the free volume is proportional to the amount of hardener in the system, as observed by Soles and co-workers[29,31], Grave and co-workers[30], and Carfagna and co-workers[32].

5. Conclusions

Figure 4. Characteristic curve obtained by applying the Berens‑Hopfenberg model to the experimental mass gain vs. time data points.

saturation[31]. This abnormal behavior is related to polymer degradation caused by water absorption[15,33]. The degradation occurs by formation of hydrogen bonds between the water molecules and polar groups present in the polymer chain, causing rupture of the initial network[15]. The water absorption not only causes plasticization of the resin, but also causes a change in the stress state, what favors the formation of cracks by swelling. These phenomena contribute to an increase in the variation of the internal structure of the material and cause an increasing weight gain (M%) close to the equilibrium level, so that this level is never achieved[33]. The curves obtained with the Berens-Hopfenberg model, Table 3, showed similar results for all phr ratios, with little variation in the diffusivity (D) values and continued increase of the saturation values related to Fickian diffusion (M∞,F) and also due to polymer swelling (M∞,R). The final behavior was similar to the one obtained with the Jacobs‑Jones models (Table 2). Polímeros, 26(1), 30-37, 2016

The use of different diffusion models allowed characterizing the absorption behavior of the DGEBA/TETA system with different hardener/epoxy ratios. The models used showed excellent convergence, with the exception of the Carter and Kibler model. Weight gain due to water absorption increased with increasing the hardener content, while diffusivity followed the opposite behavior. The trend observed for these two parameters could be explained by the increase of the free volume within the material, and by interactions caused by chain polarity and topology. The behavior observed for all systems, i.e., for all hardener/epoxy monomer ratios, followed a non-Fickian trend. The best fit was obtained when the Jacobs-Jones models were used, characterizing the presence of two phases in the material. The denser phase, i.e., the one with a higher crosslink density, is present in major proportion, and a less dense phase, with a lesser number of crosslinks is present in a smaller proportion. For some formulations, especially the one with phr 17 when the Jacobs-Jones model was used, or the ones with phr 9, 11 and 17 when the modified Jacobs-Jones model is used, the data points evidenced the occurrence of degradation due to water absorption.

6. Acknowledgements The authors acknowledge the grants from the Brazilian funding agency CNPq. 35


Pereira, A. A. C., & d’ Almeida, J. R. M.

7. References 1. Lee, H., & Neville, K. (1982). Handbook of epoxy resins. New York: McGraw-Hill. 2. Groß, A., Kollek, H., Schormann, A., & Brockmann, H. (1988). Spectroscopical contributions to the regioselectivity of nucleophilic curing reactions in epoxy resins. International Journal of Adhesion and Adhesives, 8(3), 147-158. http://dx.doi. org/10.1016/0143-7496(88)90093-0. 3. Jiang, S., Zha, S., Xia, L., & Guan, R. (2015). Synthesis and characterization of diphenylsilanediol modified epoxy resin and curing agent. Journal of Adhesion Science and Technology, 29(7), 641-656. http://dx.doi.org/10.1080/01694243.2014.10 03177. 4. Morgan, R. J., & Mones, E. T. (1987). The cure reactions, network structure, and mechanical response of diaminodiphenyl sulfone-cured tetraglycidyl 4,4’diaminodiphenyl methane epoxies. Journal of Applied Polymer Science, 33(4), 999-1020. http://dx.doi.org/10.1002/app.1987.070330401. 5. Thomas, R., Yumei, D., Yuelong, H., Le, Y., Moldenaers, P., Weimin, Y., Czigany, T., & Thomas, S. (2008). Miscibility, morphology, thermal, and mechanical properties of a DGEBA based epoxy resin toughened with a liquid rubber. Polymer, 49(1), 278-294. http://dx.doi.org/10.1016/j.polymer.2007.11.030. 6. d’Almeida, J. R. M., & Monteiro, S. N. (1996). Analysis of the fracture surface morphology of an epoxy system as a function of the resin/hardener ratio. Journal of Materials Science Letters, 15, 955-958. http://dx.doi.org/10.1007/BF00241436. 7. d’Almeida, J. R. M., & Monteiro, S. N. (1998). The influence of the amount of hardener on the tensile mechanical behavior of an epoxy system. Polymers for Advanced Technologies, 9(3), 216-221. http://dx.doi.org/10.1002/(SICI)10991581(199803)9:3<216::AID-PAT746>3.0.CO;2-S. 8. d’Almeida, J. R. M., Cella, N., Monteiro, S. N., & Miranda, L. C. M. (1998). Thermal diffusivity of an epoxy system as a function of the hardener content. Journal of Applied Polymer Science, 69(7), 1335-1341. http://dx.doi.org/10.1002/(SICI)10974628(19980815)69:7<1335::AID-APP8>3.0.CO;2-F. 9. Chiao, L., & Lyon, R. E. (1990). A fundamental approach to resin cure kinectics. Journal of Composite Materials, 24(7), 739-752. http://dx.doi.org/10.1177/002199839002400704. 10. Bouchonneau, N., Carvalho, A. R., Macêdo, A. R. L., Viana, L. U., Nascimento, A. P., Duarte, J. B. F., & Macêdo, A. R. M. (2010). Análise da absorção de água em dois polímeros expandidos: desenvolvimento do módulo de flutuabilidade de um mini-robô submarino. Polímeros: Ciência e Tecnologia, 20(3), 181-187. http://dx.doi.org/10.1590/S0104-14282010005000032. 11. Crank, J. (1975). The mathematics of diffusion (2nd ed). Oxford: Clarendon Press. 12. Shen, C. H., & Springer, G. S. (1976). Moisture absorption and desorption of composite materials. Journal of Composite Materials, 2(1), 2-20. http://dx.doi.org/10.1177/002199837601000101. 13. McKague, E. L., Reynolds, J. D., & Halkias, J. E. (1976). Moisture Diffusion in Fiber Reinforced Plastics. Journal of Engineering Materials and Technology, 98(1), 92-95. http:// dx.doi.org/10.1115/1.3443342. 14. Bortolin, A., Aouada, F. A., Longo, E., & Mattoso, L. H. C. (2012). Investigação do processo de absorção de água de hidrogéis de polissacarídeo: efeito da carga iônica, presença de sais, concentrações de monômero e polissacarídeo. Polímeros: Ciência e Tecnologia, 22(4), 311-317. http://dx.doi.org/10.1590/ S0104-14282012005000046. 15. Papanicolaou, G. C., Kosmidou, Th. V., Vatalis, A. S., & Delides, C. G. (2006). Water absorption mechanism and some anomalous effects on the mechanical and viscoelastic behavior 36

of an epoxy system. Journal of Applied Polymer Science, 99(4), 1328-1339. http://dx.doi.org/10.1002/app.22095. 16. Manfredi, L. B., Santis, H., & Vásquez, A. (2008). Influence of the addition of montmorillonite to the matrix of unidirectional glass fibre/epoxy composites on their mechanical and water absorption. Composites Part A: Applied Science and Manufacturing, 39(11), 1726-1731. http://dx.doi.org/10.1016/j. compositesa.2008.07.016. 17. Ahmad, Z., Ansell, M. P., & Smedley, D. (2011). Moisture absorption characteristics of epoxy based adhesive reinforced with CTBN and ceramic particles for bonded-in timber connection: Fickian or Non-Fickian behavior. IOP Conference Series. Materials Science and Engineering, 17, 012011. http:// dx.doi.org/10.1088/1757-899X/17/1/012011. 18. Carter, H. G., & Kibler, K. G. (1978). Langmuir-Type model for anomalous moisture diffusion in composite resins. Journal of Composite Materials, 12(2), 118-131. http://dx.doi. org/10.1177/002199837801200201. 19. Bonniau, P., & Bunsell, A. R. (1981). A comparative study of water absorption theories applied to glass epoxy composites. Journal of Composite Materials, 15(3), 272-293. http://dx.doi. org/10.1177/002199838101500306. 20. Bonniau, P., & Bunsell, A. R. (1981). Water absorption by glass fibre reinforced epoxy resin. In I. H. Marshall (Ed.), Composite structures (pp. 92-105). Amsterdam: Springer. http://dx.doi. org/10.1007/978-94-009-8120-1_7. 21. Jacobs, P. M., & Jones, F. R. (1989). Diffusion of moisture into two-phase polymers. Journal of Materials Science, 24(7), 2331-2336. http://dx.doi.org/10.1007/BF01174492. 22. Jacobs, P. M., & Jones, F. R. (1989). Diffusion of moisture into two-phase polymers. Journal of Materials Science, 24(7), 2343-2347. http://dx.doi.org/10.1007/BF01174494. 23. Berens, A. R., & Hopfenberg, H. B. (1978). Diffusion and relaxation in glassy polymer powders: 2. Separation of diffusion and relaxation parameters. Polymer, 19(5), 489-496. http:// dx.doi.org/10.1016/0032-3861(78)90269-0. 24. Pritchard, G., & Speake, S. D. (1987). The use of water absorption kinetic data to predict laminate property changes. Composites, 18(3), 227-232. http://dx.doi.org/10.1016/00104361(87)90412-5. 25. Maggana, C., & Pissis, P. (1999). Water sorption and diffusion studies in na epoxy resin system. Journal of Polymer Science Part B, 37(11), 1165-1182. http://dx.doi.org/10.1002/(SICI)10990488(19990601)37:11<1165::AID-POLB11>3.0.CO;2-E. 26. Popineau, S., Rondeau-Mouro, C., Sulpice-Gaillet, C., & Shanahan, M. E. R. (2005). Free/bound water absorption in an epoxy adhesive. Polymer, 46(24), 10733-10740. http:// dx.doi.org/10.1016/j.polymer.2005.09.008. 27. Abdelkader, A. F., & White, J. R. (2005). Water absorption in epoxy resins: The effects of the crosslinking agent and curing temperature. Journal of Applied Polymer Science, 98(6), 25442549. http://dx.doi.org/10.1002/app.22400. 28. Soles, C. L., Chang, F. T., Gidley, D. W., & Yee, A. F. (2000). Contributions of the nanovoid structure to the kinectics of moisture transport in epoxy resins. Journal of Polymer Science Part B, 38(5), 776-791. http://dx.doi.org/10.1002/(SICI)10990488(20000301)38:5<776::AID-POLB15>3.0.CO;2-A. 29. Soles, C. L., Chang, F. T., Bolan, B. A., Hristov, H. A., Gidley, D. W., & Yee, A. F. (1998). Contributions of the nanovoid structure to the moisture absorption properties of epoxy resins. Journal of Polymer Science Part B, 36(17), 3035-3048. http:// dx.doi.org/10.1002/(SICI)1099-0488(199812)36:17<3035::AIDPOLB4>3.0.CO;2-Y. 30. Grave, C., Mcewan, I., & Pethrick, R. A. (1998). Influence of stoichiometric ratio on water absorption in epoxy resins. Journal of Applied Polymer Science, 69(12), 2369-2376. http://dx.doi. Polímeros, 26(1), 30-37, 2016


Effect of the hardener to epoxy monomer ratio on the water absorption behavior of the DGEBA/TETA epoxy system org/10.1002/(SICI)1097-4628(19980919)69:12<2369::AIDAPP8>3.0.CO;2-6. 31. Soles, C. L., & Yee, A. F. (2000). A discussion of the molecular mechanisms of moisture transport in epoxy resins. Journal of Polymer Science Part B, 38(5), 792-802. http://dx.doi. org/10.1002/(SICI)1099-0488(20000301)38:5<792::AIDPOLB16>3.0.CO;2-H. 32. Carfagna, C., Apicella, A., & Nicolais, L. (1982). The effect of the prepolymer composition of amino-hardened epoxy resins on the water sorption behavior and plasticization. Journal

PolĂ­meros, 26(1), 30-37, 2016

of Applied Polymer Science, 27(1), 105-112. http://dx.doi. org/10.1002/app.1982.070270112. 33. Hahn, H. T. (1987). Hygrothermal damage in graphite/epoxy laminates. Journal of Engineering Materials and Technology, 109(1), 3-11. http://dx.doi.org/10.1115/1.3225930. Received: Mar. 06, 2015 Revised: Aug. 15, 2015 Accepted: Aug. 31, 2015

37


http://dx.doi.org/10.1590/0104-1428.1808

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

PF/CLAY hybrid materials: a simple method to modulate the optical properties Marcio Chao Chen Em1, Camila Gouveia Barbosa1, Laura Oliveira Péres1* and Roselena Faez2 Laboratory of Hybrid Materials, Institute of Environmental, Chemical and Pharmaceutical Sciences, Universidade Federal de São Paulo – UNIFESP, Diadema, SP, Brazil 2 Laboratory of Polymeric and Biosorbent Materials, Department of Natural Science, Mathematics and Education, Universidade Federal de São Carlos – UFSCar, Araras, SP, Brazil

1

*lauraoperes@gmail.com

Abstract The aim of this work was modulate the emission properties and improve thermal stability of a conjugated polymer incorporated into an inorganic matrix. Hybrid material was prepared based on poly(9,9-dioctylfluorene-co-phenylene (PF) and montmorillonite (Na+Mt) clay using wet impregnation of 10, 30 and 50 wt.% of PF into Na+Mt and Na+Mt intercalated with ammonium quaternary salts (hexadecyltrimethylammonium — HDTMA) in a different proportions (OMt-1 and OMt-2). The materials were characterized by infrared and UV-Vis spectroscopy, fluorescence, X-ray diffratometry and thermogravimetry analysis. The results show that the presence of the clay alters the photoluminescent and thermal properties. Nevertheless, the degree of the clay organophilization and the clay content influences the luminescent properties due to the diverse interaction behavior between the polymer and clay. The sodium clay acted only as dispersing agent since no intercalation process occurs and the emission displacement is assigned to this behavior. In this case the PF emission displace from 402 to 395 nm. A nonlinear displacement is observed for PF/OMt-2 due the difficulties to conclude if the intercalation of the polymer occurs (379, 403 and 412 for hybrid with 10, 30 and 50%, respectively). For PF/OMt-1 a higher displacements for lower wavelength is observed due to intercalation of polymer chains and subsequent isolation in the interlamellar space, especially with material with 10 and 30% of PF in the hybrid material, whose displacement reached to 360 nm. All these results show that is possible to try to control the emission of the conjugated hybrid material changing the rate of the material. Keywords: polyfluorene, clay, hybrid materials, photoluminescence.

1. Introduction Among the luminescent conjugated polymers polyfluorene (PF) family, PF is one of the most used and developed due to its high chemical and thermal stabilities. This behavior reflects the spectral stability allowing the formation of liquid-crystalline phase and thus supports the manufacture of electroluminescent device with highly polarized light emission[1]. PF have high efficiency in the blue region and can be prepared as copolymers with others light emitting polymers[2]. Furthermore, it can be modulated to emit in the visible region ranging from blue to red[3,4]. This can occur when substituents are inserted in the polymer chain and affect the properties such as, solubility, thermal stability, conductivity, optical and electrochemical properties[3,5,6]. Also, materials can be obtained with high photoluminescence quantum efficiency and good quantum transport charges[5,7-9]. Hybrid inorganic/organic materials have been extensively studied mainly because of their intrinsic properties obtained by a synergistic effect of its components[10-17]. The combination of conjugated polymers and inorganic matrix such as layered silicates, ordered mesoporous silica and carbon nanotubes have great interest because of the potential applications in electrical and electronic sectors[18-21]. Among the layered silicates class the mineral clays presents important characteristics, besides of that the naturally occurring materials, such

38

cation exchange capacity, high surface area and capacity for adsorption/absorption properties[16,22-25]. It is also well known to the scientific community that the synergistic combination of these materials that improve the photoluminescence quantum efficiency[26-30]. Further the presence of inorganic phase can also modify the color of emission[31]. The displacement of the emission peak may be associated to the decrease of the polymer aggregation process which the optical properties of conjugated polymers are dependent of chain conformation. It is known that the isolation of conjugated polymer chains within inorganic materials plays an important role in improving the luminescent efficiency and controlling interchain transfer energy[31,32]. Different methodologies can be used to combine the polymer and inorganic host. The polymer chain can be inserted by intercalation when the inorganic matrix has lamellar structure such as natural aluminosilicates. These nanohybrids of polymer/clay have an improved properties when compared with their analogues pure polymers, such as thermal and mechanical stability, chemical resistance and performance against photodegradation[27,33]. In this context, we report the preparation of montmorilonite and PF hybrids materials and their structural (FTIR and XRD), thermal (TG) and luminescence characterization.

Polímeros, 26(1), 38-43, 2016


PF/CLAY hybrid materials: a simple method to modulate the optical properties The drive for the use of lamellar structure as montmorillonite clay was related to its good thermal and chemical stability and the possibility to isolate the polymer chain.

2. Experimental Section 2.1 Polymer synthesis Poly(9,9-dioctylfluorene)-co-phenylene (PF) was synthesized as describe in the literature[32,33] using 9,9-dioctylfluorene2,7-dibromofluorene, 1,4-phenylenebisboronic acid and tetrakis-(triphenylphosphine) palladium (P(Ph3)4Pd) as a catalyst.

2.2 Organophilization of montmorilonitte Sodium montmorillonite clay, (Na+Mt) (Bentonit União Nordeste, Brazil) with a cation exchange capacity (CEC) of 85 meq/100 g of clay was organophilized with hexadecyltrimethylammonium bromide (HDTMAB, Aldrich) according to the Fontana et al.[34] The mineral clay was dispersed in an aqueous HDTMA solution using weight ratios (HDTMA:Na+Mt) of 1:1 and 1:2 based on the CEC of the clay. The organomodified clays (OMt) were respectively named as OMt-1.0 and OMt -2.0.

3. Results and Discussion As previous pointed out, PF/Na+Mt and PF/OMt composites were prepared by dispersing PF and both clays in a solvent. After that solid materials were separated, dried and used for structural (FTIR and XRD), thermal (thermogravimetry) and absorption-emission (UV-Vis and photoluminescence) characterization. FTIR spectra of PF and its hybrids prepared with 50 wt.% are shown in Figure 1, and the Table 1 summarizes the main bands and their assignments. The characteristic bands of the polyfluorene are verified mainly at 1599 cm–1 to C=C stretching vibrations, 1460 cm–1 to CH2 bending of fluorine group and 812 and 760 cm–1 assigned to C-H stretching vibrations. The bands of boronic group (1320 and 1350 cm–1), presented in the started materials, cannot be shown, indicating the formation of the polymer. In the hybrids it is possible to observe the silicate related absorption bands at 1030 cm–1 for Si-O-Si stretching and 528 and 470 cm–1 for Si-O stretching and bending, respectively. The presence of the bands at 2925 and 2854 cm–1 (aliphatic CH stretching) arising PF and clay can be seen. But, the presence of the PF, in the hybrid, can be confirmed through the bands at 1599 cm–1 (aromatic ring stretching) and 1300 cm–1 (CH

2.3 Preparation of PF/clay In a Erlenmeyer, equipped with a magnetic stirring, 0.5 g of the clays were first dispersed in 18 mL of ethanol (Na+Mt) and toluene (OMt-1 and OMt-2) and 0.05 g of PF (for 10% in mass of polymer) was dissolved in toluene. After complete dispersion the solutions were mixed and kept closed under magnetic stirring for 48 hours. Subsequently, the solvent was removed and the material was dried. PF-clay hybrids were prepared with 10, 30 and 50% in mass of polymer and respectively named as Na+Mt/PF(10), Na+Mt/PF(30), Na+Mt/PF(50), OMt-1/PF(10), OMt-1/PF(30), OMt-1/PF(50), OMt-2/PF(10), OMt-2/PF(30) and OMt-2/PF(50).

2.4 Characterization UV absorption spectra were collected on a Varian Cary 50 spectrophotometer. The materials were dispersed in spectroscopic chloroform (1 × 10–5 mol/L), with the solution being poured into a 10 mm square quartz cell. The absorption spectrum was collected in the range λ= 200‑600 nm. Photoluminescence (PL) spectra were taken with a Varian Eclipse fluorescence spectrophotometer. Powder sample was pressed into two plates of glass and excited at 315 nm. The spectra of the solution were performed in spectroscopic chloroform (1 × 10–5 mol/L), with excitation at the wavelength of the maximum absorption according to the UV-Vis spectra. TG curves were obtained in a thermogravimetric analyzer TG-60/60H (Shimadzu) from 50 °C to 900 °C at a heating rate of 10 °C.min–1 under nitrogen atmosphere (200 mL/min), sample weight 10 mg. Infrared spectra were recorded in the range 4000-400 cm–1 in KBr pellets, using a BOMEM MB-100 spectrometer. X ray diffraction (XRD) patterns of powdered samples were recorded on a Rigaku diffractometer model Miniflex using Cu-Kα radiation (1.541 Å, 30 kV and 15 mA). Polímeros, 26(1), 38-43, 2016

Figure 1. FTIR spectra of (a) PF and hybrids prepared with 50 wt.% of (b) Na+Mt; (c) OMt-1 and (d) OMt-2.

Table 1. Position and respective assignment of the main bands present at the FTIR spectra of PF and hybrids materials. Band Position (wavenumber cm–1) 528 and 470 812 and 760 1030 1300 1460 1599 2925 and 2854

Assignment Si-O stretching and bending C-H stretching vibrations Si-O-Si stretching CH in-plane deformation CH2 bending of fluorine group C=C stretching vibrations aliphatic CH stretching

39


Em, M. C. C., Barbosa, C. G., Péres, L. O., & Faez, R. in-plane deformation), as can be seen in the dash line in the graph. Throughout the spectral region no displacement and/ or appearance of new bands are observed in the composites suggesting no chemical interaction between the components. Similar results were observed for other proportions. XRD difractograms of clays and its hybrids in a proportion of 50 wt.% are shown in Figure 2. The organomodification of the Na+Mt with HDTMA+ was successfully realized. The interlayer space was increased as HDTMA+ was inserted in the basal area as observed by an increase in interlayer

Figure 2. X-ray diffraction patterns of (a) PF/ Na+Mt; (b) Na+Mt; (c) PF/OMt-1; (d) OMt-1; (e) PF/OMt-2; (f) OMt-2 and (g) PF.

space from 1.4 to 2.2 and 3.6 nm for Na+Mt to OMt-1 and OMt-2, respectively[34,35]. Figure 2 also compare the XRD patterns of the clays with its hybrids. From Na+Mt and Na+Mt/PF(50) no changes in the basal spacing is observed suggesting the Na+Mt only acted as a dispersed phase to dilute and de-aggregate the polymer chains. Comparing OMt-1 and OMt-1/PF(50) it is observed a decrease on the peak associate with basal plane d001 increasing the interlayer space from 2.18 to 4.20 nm. These results strongly suggest the intercalation of the polymer chain and also corroborated to the emission spectra of the hybrids, as discussed later. For OMt-2/PF(50) a small dislocations in the d001 of OMt-2 is observed. However, it is observed the appearance of new peak at 2θ equal 3.4°. The d001 peak dislocation and the formation of a new phase (arrow in the Figure 2f) could suggest the presence of the polyfluorene in the clay. Thermal behavior of PF, clay and hybrids was evaluated by thermogravimetry analysis in nitrogen atmosphere and Figure 3 show the TG and DTG curves. For PF only one stage of weight loss is observed (Ti = 284 °C) addressed to the main chain degradation[36]. The PF/Na+Mt(50) present two stage of thermodecomposition. The first one assigned to the water loss and the second one (from 230 to 410 °C) to the PF degradation (50% of weigh loss). PF/OMt-1(50) and PF/OMt-2(50) shows three stage of weight loss. The first one assigned mainly to the quaternary ammonium ion degradation and the second and third stages, from 220 to 450 °C, due to the ammonium ion and PF degradation. The DTG curves show better the event separation, Figure 4. However, for higher amount of quaternary salt lower the initial temperature of degradation since the ammonium salt has lower stability. For others compositions the results are similar (not shown here). Figure 4 shows the absorption and emission curves in chloroform solution of PF and hybrids prepared with Na+Mt, OMt-1 and OMt-2. It is not observed displacement and/or appearance of new bands on hybrids spectra compared to the bulk polymer. It is possible to realize that, the excitation of the electrons from the valence band to the conduction band, ie, the beginning of the π−π* transition, occurs at 350 nm (3.5 eV). This band gap energy, corresponding to the traditional band gap range of semiconductor materials,

Figure 3. TG and DTG curves of (a) pure PF and PF with 50 wt.% of (b) Na+Mt; (c) OMt-1 and (d) OMt-2. 40

Polímeros, 26(1), 38-43, 2016


PF/CLAY hybrid materials: a simple method to modulate the optical properties shows the conductive properties of this polymer. In the emission spectra the maxima are found at 360 and 377 nm. For hybrids there is an inversion of the bands intensities

Figure 4. Absorption and emission spectra in chloroform of (a) pure PF and hybrids of PF prepared with 50 wt. % of (b) Na+Mt; (c) OMt-1 and (d) OMt-2.

compared to the pure polymer. This behavior may be related to the vibronic structure which is associated to the C=C stretching mode band (at 1599 cm–1 in the infrared spectrum)[37]. Normally, the absorption and emission spectra are specular. This symmetry is due to the transitions involved, i.e., for emission spectra the transition occurs of an excited state n to the fundamental and for absorption spectra the process is reversed. In real polymer systems is difficult to separate those different processes due to the widths of bands and thus the specularity normally is not observed. Usually the spectra are broad due to interchain processes associated with the effects of homogeneity of the film, conformational disorder and multiple processes of decays. Moreover, the energy transition π−π* depends on the conjugation length, the distribution of absorption energy expansion and the vibronic structure of a particular segment. Instead, the emission spectra exploiting only sites of lower energy. Accordingly, two possibilities can explain the obtained results in solution, that is, no displacement to the pure polymer; the portion of polymer adsorbed on the clays surface, out of the lamellae, have the main contribution to the emission of the material, or the lability makes the polymer comes out of the lamellae. The same behaviors are observed for others hybrids. Figure 5 shows the solid state emission spectra for PF and hybrid materials. A shift of the maximum peak for solid

Figure 5. Solid state spectra of pure PF and hybrids of PF with (a) Na+Mt; (b) OMt-1 and (c) OMt-2. Polímeros, 26(1), 38-43, 2016

41


Em, M. C. C., Barbosa, C. G., Péres, L. O., & Faez, R. spectrum of PF was verified compared the one obtained in solution (from 377 to 402 nm). For Na+Mt/PF, Figure 5a, a single broadband at 395 nm is observed regardless the amount of PF and a blue shift from 402 to 395 nm is verified. This result is in agreement with the XDR suggesting the polymer is not into the clay, but this can act as a dispersed phase to dilute the polymer chain. For OMt-1/PF, Figure 5b, there is also a blue shift of the maximum peak. However, the amount of polymer affects the band displacement. For OMt‑1/PF(50) a displacement from 402 to 380 nm was observed whereas for OMt-1/PF(10) and OMt-1/PF(30) a shift from 402 to 360 nm. For hybrids prepared OMt-2 the change in the spectra compare to PF was even greater. Initially no linearity in the behavior was observed or just blue shift when clay was added. Rationalizing and analyzing only the hybrids a red shift is detected increasing the polyfluorene and a maximum at 379, 403 and 412 nm for the hybrids with 10, 30 and 50 wt.%, respectively is found. The above results can be explained based on the aggregation and de-aggregation phenomenon. According to Scherf and List[38], the agglomeration process of the particles shifts the emission spectrum to longer wavelengths (red). This behavior is observed in the solution and solid state spectra of bulk polyfluorene. As a consequence, as more isolated species the higher is the frequency observed[39]. More isolated species (more de-aggregated) are found for hybrids prepared with OMt-1 at low percentage (10 and 30 wt.%). In fact, the process of isolation-association (aggregation-de-aggregation) is not the only phenomenon observed in the hybrids. An important statement is related to the folding process of the polymer and the mechanism of polymer intercalation in the clays. Ramachandran et al.[28] suggests that the material can be fully, partially or non-intercalated. The larger the quantity of material is sandwiched within the lamellae, the higher the red shift, due to the higher amount of aggregated material. Moreover, when the polymer chains are interspersed, the effective conjugation length increases due to the planarization of the chain. In this case the transfer of the Förster energy is difficult and, instead, the intrachain energy transfer can become dominant. Therefore, the higher-energy excitons migrate to the lowest energy state along the chain, resulting in a red shift of the spectrum[10]. The intercalation process can also provide an additional environment to change the emission color by changing the effective conjugation length. Finally, for OMt2/PF(50), three factors influence the observed shift: (i) larger intercalation of polymeric material, (ii) increase of the chains association, (iii) increase the intrachain energy transfer. 10 wt.%.

4. Conclusions Through simple technique, via wet impregnation, was possible to prepare luminescent materials consisting of polyfluorene and montmorillonite clay. Modulated emission values were obtained depending on the clay modification and the amount of polymeric material. The results suggest no chemical interaction between the components, but the X-ray diffraction strongly suggests the intercalation of the polymer chain and also corroborated to the emission spectra of the hybrids. Based on the results we can indicate this composite to be used as active material in an optoelectronic devices. 42

5. Acknowledgements The authors thank FAPESP (07/50742-2), CNPq, Capes-nBioNet and INEO-INCT (National Institute of Science and Technology for Organic Electronics) for financial support. The authors also thank Prof. Dr. Marcos Augusto Bizeto and Prof. Dr. Tereza da Silva Martins for DRX analysis.

6. References 1. Bernius, M. T., Inbasekaran, M., O’Brien, J., & Wu, W. S. (2000). Progress with light-emitting polymers. Advanced Materials, 12(23), 1737-1750. http://dx.doi.org/10.1002/15214095(200012)12:23<1737::AID-ADMA1737>3.0.CO;2-N. 2. Cirpan, A., Ding, L., & Karasz, F. E. (2005). Optical and electroluminescent properties of polyfluorene copolymers and their blends. Polymer, 46(3), 811-817. http://dx.doi. org/10.1016/j.polymer.2004.11.107. 3. Akcelrud, L. (2003). Electroluminescent polymers. Progress in Polymer Science, 28(6), 875-962. http://dx.doi.org/10.1016/ S0079-6700(02)00140-5. 4. List, E. J. W., Guentner, R., Scanducci de Freitas, P., & Scherf, U. (2002). The effect of keto defect sites on the emission properties of polyfluorene-type materials. Advanced Materials, 14(5), 374-378. http://dx.doi.org/10.1002/15214095(20020304)14:5<374::AID-ADMA374>3.0.CO;2-U. 5. Oliveira, H. P. M., Cossiello, R. F., Atvars, T. D. Z., & Akcelrud, L. (2006). Dispositivos poliméricos eletroluminescentes. Quimica Nova, 29(2), 277-286. http://dx.doi.org/10.1590/ S0100-40422006000200019. 6. Kaya, I., Yıldırım, M., Aydın, A., & Senol, D. (2010). Synthesis and characterization of fluorescent graft fluorene-co-polyphenol derivatives: the effect of substituent on solubility, thermal stability, conductivity, optical and electrochemical properties. Reactive & Functional Polymers, 70(10), 815-826. http:// dx.doi.org/10.1016/j.reactfunctpolym.2010.07.013. 7. Shakutsui, M., Matsuura, H., & Fujita, K. (2009). Improved efficiency of polymer light-emitting diodes by inserting a hole transport layer formed without thermal treatment above glass transition temperature. Organic Electronics, 10(5), 834-842. http://dx.doi.org/10.1016/j.orgel.2009.04.004. 8. Stéphan, O., Collomb, V., Vial, J.-C., & Armand, M. (2000). Blue-green light-emitting diodes and electrochemical cells based on a copolymer derived from fluorine. Synthetic Metals, 113(3), 257-262. http://dx.doi.org/10.1016/S0379-6779(00)00214-9. 9. Zhao, W., Cao, T., & White, J. M. (2004). On the origin of green emission in polyfluorene polymers: the roles of thermal oxidation degradation and crosslinking. Advanced Functional Materials, 14(8), 783-790. http://dx.doi.org/10.1002/adfm.200305173. 10. Lee, T.-W., Park, O. O., Kim, J.-J., Hong, J.-M., & Kim, Y. C. (2001). Efficient photoluminescence and electroluminescence from environmentally stable polymer/clay nanocomposites. Chemistry of Materials, 13(6), 2217-2222. http://dx.doi. org/10.1021/cm010201h. 11. Barbosa, R., Araújo, E. M., Oliveira, A. D., & Melo, T. J. A. (2006). Efeito de sais quaternários de amônio na organofilização de uma argila bentonita nacional. Cerâmica, 52(324), 264-268. http://dx.doi.org/10.1590/S0366-69132006000400009. 12. Colvin, V. L., Schlamp, M. C., & Alivisatos, A. P. (1994). Lightemitting diodes made from cadmium selenide nanocrystals and a semiconducting polymer. Nature, 370(6488), 354-357. http://dx.doi.org/10.1038/370354a0. 13. Hagrman, P. J., Hagrman, D., & Zubieta, J. (1999). Organic inorganic hybrid materials: from “simple” coordination polymers to organodiamine-templated molybdenum oxides. Angewandte Polímeros, 26(1), 38-43, 2016


PF/CLAY hybrid materials: a simple method to modulate the optical properties Chemie International Edition, 38(18), 2638-2684. http://dx.doi. org/10.1002/(SICI)1521-3773(19990917)38:18<2638::AIDANIE2638>3.0.CO;2-4. PMid:10508356. 14. Santos, M. A., Mattoso, L. H. C., Defácio, R., & Jamshid, A. (2001). Compósitos de borracha natural com compostos condutivos à base de negro de fumo e polímero condutor. Polímeros: Ciência e Tecnologia, 11(3), 126-134. http://dx.doi. org/10.1590/S0104-14282001000300012. 15. Kim, J., Kim, B., Anand, C., Mano, A., Zaidi, J. S. M., Ariga, K., You, J., Vinu, A., & Kim, E. (2015). A single-step synthesis of electroactive mesoporous ProDOT-silica structures. Angewandte Chemie International Edition, 54(29), 8407-8410. http://dx.doi. org/10.1002/anie.201502498. PMid:26037244. 16. Nakao, A., & Fujiki, M. (2015). Visualizing spontaneous physisorption of non-charged π-conjugated polymers onto neutral surfaces of spherical silica in nonpolar solvents. Polymer Journal, 47(6), 434-442. http://dx.doi.org/10.1038/ pj.2015.14. 17. Joshi, P. B., & Zhang, P. (2015). Facile capture of conjugated polymer nanodots in silica nanoparticles to facilitate surface modification. Journal of Materials Science, 50(10), 3597-3603. http://dx.doi.org/10.1007/s10853-015-8920-5. 18. Paiva, L. B., Morales, A. R., & Díaz, F. R. V. (2008). Argilas organofílicas: características, metodologias de preparação, compostos de intercalação e técnicas de caracterização. Cerâmica, 54(330), 213-226. http://dx.doi.org/10.1590/S036669132008000200012. 19. De Barros, A., Ferreira, M., Constantino, C. J. L., & Ferreira, M. (2014). Nanocomposites based on LbL films of polyaniline and sodium montmorillonite clay. Synthetic Metals, 197, 119125. http://dx.doi.org/10.1016/j.synthmet.2014.09.001. 20. Fatnassi, M., & Es-Souni, M. (2015). Nanoscale phase separation in laponite–polypyrrole nanocomposites. Application to electrodes for energy storage. Royal Society of Chemistry, 5(28), 21550-21557. http://dx.doi.org/10.1039/C4RA16540C. 21. Feng, L., Sha, J., He, Y., Chen, S., Liu, B., Zhang, H., & Lü, C. (2015). Conjugated polymer and spirolactam rhodamine-B derivative co-functionalized mesoporous silica nanoparticles as the scaffold for the FRET-based ratiometric sensing of mercury (II) ions. Microporous and Mesoporous Materials, 208, 113119. http://dx.doi.org/10.1016/j.micromeso.2015.01.039. 22. Silva, A. A., Valenzuela-Diaz, F. R., Martins, G. S. V., & Rodrigues, M. G. F. (2007). Preparation of organophilic clays using different concentrations of quaternary ammonium salt. Cerâmica, 53(328), 417-422. http://dx.doi.org/10.1590/S036669132007000400013. 23. Park, J. H., Lim, Y. T., Park, O. O., Yu, J.-W., Kim, J. K., & Kim, Y. C. (2004). Enhanced quantum efficiency in blue-emitting polymer/dielectric nanolayers nanocomposite light-emitting devices. Materials Science and Engineering C, 24(1-2), 75-78. http://dx.doi.org/10.1016/j.msec.2003.09.039. 24. Zheng, M., Ding, L., Lin, Z., & Karasz, F. E. (2002). Synthesis and Characterization of Fluorenediylvinylene and Thiophenediylvinylene-Containing Terphenylene-Based Copolymers. Macromolecules, 35(27), 9939-9946. http:// dx.doi.org/10.1021/ma020533n. 25. Ramôa, S. D. A., Merlini, C., Barra, G. M. O., & Soares, B. G. (2014). Obtenção de nanocompósitos condutores de montmorilonita/polipirrol: Efeito da incorporação do surfactante na estrutura e propriedades. Polímeros: Ciência e Tecnologia, 24(ESP), 57-62. http://dx.doi.org/10.4322/polimeros.2014.051. 26. Lee, T. W., Park, O. O., Yoon, J., & Kim, J. (2001). Enhanced quantum efficiency in polymer/layered silicate nanocomposite light emitting devices. Synthetic Metals, 121(1), 1737-1738. http://dx.doi.org/10.1016/S0379-6779(00)01493-4.

Polímeros, 26(1), 38-43, 2016

27. Jing, C., Chen, L., Shi, Y., & Jin, X. (2005). Synthesis and characterization of exfoliated MEHPPV/clay nanocomposites by in situ polymerization. European Polymer Journal, 41(10), 2388-2394. http://dx.doi.org/10.1016/j.eurpolymj.2005.05.007. 28. Ramachandran, G., Simon, G. P., Cheng, Y. B., & Dai, L. (2005). Control of fluorescence emission color of benzo 15-crown-5 ether substituted oligo phenylene vinylene-ceramic nanocomposites. Polymer, 46(18), 7176-7184. http://dx.doi. org/10.1016/j.polymer.2005.05.087. 29. Evans, R. C., Macedo, A. G., Pradhan, S., Scherf, U., Carlos, L. D., & Burrows, H. D. (2010). Fluorene based conjugated polyelectrolyte/silica nanocomposites: chatge-mediated phase aggregation at the organic-inorganic interfece. Advanced Materials, 22(28), 3032-3037. http://dx.doi.org/10.1002/ adma.200904377. PMid:20535734. 30. Winkler, B., Dai, L., & Mau, A. W.-H. (1999). Organicinorganic hybrid light-emitting composites: poly(p-phenylene vinylene) intercalated clay nanoparticles. Journal of Materials Science Letters, 18(19), 1539-1541. http://dx.doi. org/10.1023/A:1006610926414. 31. Lee, H.-C., Lee, T.-W., Lim, Y. T., & Park, O. O. (2002). Improved environmental stability in poly(p-phenylene vinylene)/ layered silicate nanocomposite. Applied Clay Science, 21(5-6), 287-293. http://dx.doi.org/10.1016/S0169-1317(02)00090-X. 32. Santos, T. C. F., Peres, L. O., Wang, S. H., Oliveira, O. N. Jr, & Caseli, L. (2010). Mixing alternating copolymers containing fluorenyl groups with phospholipids to obtain Langmuir and Langmuir-Blodgett films. Langmuir, 26(8), 5869-5875. http:// dx.doi.org/10.1021/la9038107. PMid:19921831. 33. Péres, L. O., Errien, N., Faulques, E., Athalin, H., Lefrant, S., Massuyeau, F., Wéry, J., Froyer, G., & Wang, S. H. (2007). Synthesis and characterization of a new alternating copolymer containing quaterphenyl and fluorenyl groups. Polymer, 48(1), 98-104. http://dx.doi.org/10.1016/j.polymer.2006.10.038. 34. Fontana, J. P., Camilo, F. F., Bizeto, M. A., & Faez, R. (2013). Evaluation of the role of an ionic liquid as organophilization agent into montmorillonite for NBR rubber nanocomposite production. Applied Clay Science, 83-84, 203-209. http:// dx.doi.org/10.1016/j.clay.2013.09.002. 35. Bergaya, F., Theng, B. K. G., & Lagaly, G. (2006). Handbook of Clay Science (Vol. 1, Developments in Clay Science). Amsterdam: Elsevier. 36. Xia, C., & Advincula, R. C. (2001). Decreased aggregation phenomena in polyfluorenes by introducing carbazole copolymer. Macromolecules, 34(17), 5854-5859. http://dx.doi.org/10.1021/ ma002036h. 37. Ranger, M., Rondeau, D., & Leclerc, M. (1997). New WellDefined Poly(2,7-fluorene) Derivatives: Photoluminescence and Base Doping. Macromolecules, 30(25), 7686-7691. http:// dx.doi.org/10.1021/ma970920a. 38. Scherf, U., & List, E. J. W. (2002). Semiconducting polyfluorenes: towards reliable structure: property relationships. Advanced Materials, 14(7), 477-487. http://dx.doi.org/10.1002/15214095(20020404)14:7<477::AID-ADMA477>3.0.CO;2-9. 39. Cassemiro, S. M., Thomazi, F., Roman, L. S., Marletta, A., & Akcelrud, L. (2009). Effect of conjugation length on photophysical properties of a conjugated–non-conjugated multiblock copolymer. Synthetic Metals, 159(19-20), 19751982. http://dx.doi.org/10.1016/j.synthmet.2009.07.004. Received: June 16, 2014 Revised: June 29, 2015 Accepted: Aug. 31, 2015

43


http://dx.doi.org/10.1590/0104-1428.2100

R R R R R R R R R R R R R R

Thermal and catalytic pyrolysis of plastic waste Débora Almeida1 and Maria de Fátima Marques1* Instituto de Macromoléculas Eloisa Mano, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brasil

1

*fmarques@ima.ufrj.br

Abstract The amount of plastic waste is growing every year and with that comes an environmental concern regarding this problem. Pyrolysis as a tertiary recycling process is presented as a solution. Pyrolysis can be thermal or catalytical and can be performed under different experimental conditions. These conditions affect the type and amount of product obtained. With the pyrolysis process, products can be obtained with high added value, such as fuel oils and feedstock for new products. Zeolites can be used as catalysts in catalytic pyrolysis and influence the final products obtained. Keywords: catalytic pyrolysis, fuel oils, thermal pyrolysis, zeolites.

1. Introduction Plastics are materials that offer a fundamental contribution to our society, due to its versatility and relatively low cost. As a result of this contribution, a large amount of plastic waste is generated due to the increase in its production each year. This increase in the amount of waste does cause some environmental problems, since plastics do not degrade quickly and can remain in the environment for a long time[1-5]. A large part of this waste is disposed of in landfills or is incinerated[6,7]. However, the plastic waste are bulkier than other organic residues and thus occupy massive space in landfills and therefore the proper disposal and incineration have high costs. Furthermore, incineration of these waste plastics results in environmental problems due to increased emission of harmful compounds[2,6-8]. It is necessary for more sustainable solutions that incineration and disposal in landfills are researched and developed[4]. Thus, much research in the area of recycling and reuse of these post-consumed polymers have been carried out in order to produce raw materials and energy[1,3,7]. The various types of recycling are good options to control the increase of plastic waste, because they are environmentally friendly when compared with incineration and disposal in landfills. In fact, from recycling it is possible to recover raw materials, energy and fuel while minimizing the consumption of natural resources and raw materials. When these products and energy are recovered, the environmental impacts of industrial activity are reduced[1,3,9,10]. Municipal waste plastics are heterogeneous, unlike industrial. For homogeneous plastic waste, the repelletization and remoulding can be a simple and effective means of recycling. However, when these wastes are heterogeneous and consist of mixtures of resins, they are unsuitable for such recovery. In this case, other forms of recycling[11] are necessary. Each recycling method provides a number of advantages that make them beneficial for local and specific applications[12]. Appropriate treatment of plastic waste is an important question for waste management, due to energy, environmental, economic, and political[11] aspects.

44

The plastics recycling methods, in accordance with ASTM D5033-00, are divided into four types according to the final result, one of them being the tertiary or chemical recycling. In this type of recycling chemical degradation leads to production of liquid fuels and chemicals with high added value from waste plastic fragments or segregated[2,8,13,14]. One of the tertiary recycling methods is pyrolysis. This process can be thermal or catalytic and is a promising alternative that allows the conversion of polymers into gas and liquid hydrocarbons[4,15,16]. Pyrolysis is a process with relatively low cost from which a wide distribution of products can be obtained. In the process of pyrolysis, where heating occurs in the absence of oxygen, the organic compounds are decomposed generating gaseous and liquid products, which can be used as fuels and / or sources of chemicals. Meanwhile, the inorganic material, free of organic matter, remains unchanged under the solid fraction and can be recycled later[17]. The thermal pyrolysis requires high temperatures, which often results in products with low quality, making this process unfeasible. This occurs because the uncatalyzed thermal degradation gives rise to low molecular weight substances, however in a very wide range of products[13,15,16]. This method can be improved by the addition of catalysts, which will reduce the temperature and reaction time and allow the production of hydrocarbons with a higher added value, such as fuel oils and petrochemical feedstocks[4,11,18-21]. That is, the use of catalysts gives an added value to the pyrolysis and cracking efficiency of these catalysts depends both on its chemical and physical characteristics. These particular properties, promote the breaking of C-C bonds and determine the length of the chains of the products obtained[17,22]. For Brazilian cities, the percentage of high and low density polyethylene (HDPE and LDPE, respectively), polyethylene terephthalate (PET), Poly(vinyl chloride) (PVC) and polypropylene (PP) found in municipal solid wastes are 89% and the other polymers account for the other 11%[13].

Polímeros, 26(1), 44-51, 2016


Thermal and catalytic pyrolysis of plastic waste Therefore, polyolefins (PE, PP and their copolymers) are the most widely used thermoplastics for several applications and are most of the polymeric residues, that make up 60‑70% of municipal solid waste[23]. Tertiary recycling of plastic waste containing PVC releases hydrogen chloride, which causes corrosion of the pyrolysis reactor and formation of organochlorine compounds[23]. The presence of chlorine is very harmful for use as fuel in the pyrolysis liquid products obtained[24]. Although plastic waste may be considered economical sources of chemicals and energy, recycling of mixed plastic waste containing PVC not only result in the formation of volatile organic compounds in products, also in the emission of pollution when they are applied[23]. Moreover, PET may be mechanically recycled obtaining fibers for carpets, clothes and bottles. The products obtained in this recycling are of high quality that can be compared with virgin polymer[12]. Therefore, PET and other special polymers should be removed from municipal waste by mechanical recovery, which is economically viable.

1.1 Pyrolysis The tertiary or chemical recycling includes a variety of processes that enable the generation of high value products such as fuel or chemicals[11,16,19-21,25-27]. In this method, the plastic waste is processed to produce basic petrochemical compounds, which can be used as raw material for new plastics. This process has the advantage of working with mixed and contaminated plastics[12,18,20,27]. Recently, much attention has been directed to chemical recycling, particularly the uncatalyzed thermal cracking (thermolysis), catalytic cracking and steam decomposition, as methods for producing various hydrocarbon fractions in the range of fuel, from solid waste plastics[12]. In the case of polymers, pyrolysis stands out as tertiary recycling method, however this cracking gives rise to low molecular weight substances, however unfortunately in a very wide range of products, in the case of non-catalyzed thermal decomposition[11,13,15,16,18,26]. The pyrolysis can be carried out at different temperatures, reaction times, pressures, in the presence or absence of catalysts and reactive gases. The pyrolysis process involves the breaking of bonds, and is generally endothermic and hence the supply of heat is essential to react the material[28]. In polymeric samples, the decomposition process may occur through the elimination of small molecules, chain scission (depolymerization) or random cleavage[29]. In the pyrolysis process, the sample is heated in the absence of oxygen and the organic compounds are decomposed generating gaseous and liquid products. On the other hand, the inorganic part of the sample, free from organic matter remains practically unchanged in the solid fraction enabling their separation and recovery for subsequent reuse. Therefore, the pyrolysis is an attractive alternative technique for recycling waste plastics recycling[2,8,17,24,30]. Thermal pyrolysis involves the decomposition of polymeric materials by means of temperature when it is applied under inert atmospheric conditions. This process is usually conducted at temperatures between 350 and 900 °C. Polímeros, 26(1), 44-51, 2016

In the case of polyolefins, which make up much of urban waste plastics, the process proceeds through random cleavage mechanism that generates a heterogeneous mixture of linear paraffins and olefins in a wide range of molar masses[11,18,20,21]. On the other hand, the catalyzed pyrolysis promotes these decomposition reactions at lower temperatures and shorter times, because of the presence of catalysts that assist in the process. Thus, the catalytic pyrolysis presents a number of advantages over thermal, such as lower energy consumption and product formation with narrower distribution of the number of carbon atoms, which may be directed to aromatic hydrocarbons with light and high market value[11,18-21,26]. The kinetics of degradation and the pyrolysis mechanism are still being studied and discussed. Degradation has a very complex mechanism, so adequate description of decomposing a mixture of polymers is difficult, even more so in the presence of catalysts and a process with several stages[7,30]. In order to solve this problem there are some methods based on the mass loss curve during pyrolysis. Thermogravimetric analysis (TGA) is a method that can be used to determine the loss of mass and kinetic parameters. Thermogravimetric analysis of pyrolysis involves the thermal degradation of the sample in an inert atmosphere obtaining simultaneously the weight loss values of the samples with increasing temperature at a constant heating rate[4,21,31]. Most of the techniques that are used to monitor the reactions, both for the identification of products of the gas phase and by thermogravimetric analysis, will only detect the reaction when the molecules of the products become small enough to evaporate in the gas fraction and can be observed as gas fraction or by means of mass loss of the initial sample. Is possible to follow the reaction from the beginning, since each broken link consume certain amount of energy. Thus, by measuring the heat flow into the sample during the reaction (using for example the calorimeter DSC method), it is possible to measure the rate of broken bonds occurring in the sample[4]. The reaction rates and other kinetic parameters of the degradation of the polymer are dependent on the chemical structure of these polymers. Generally the CC bonds of the polymer backbone are broken forming a higher degree of branching structures, due to the lower thermal stability of the tertiary carbon atom. Moreover, the mechanism may also be affected by contaminants. The actual reason for the differences between the rates of degradation of the macromolecules has been explained by the distortion of electron density from the degraded polymer, which depends primarily on the side group linked to the main chain of the macromolecule. For this reason, polypropylene (PP) is less stable than polyethylene (LDPE, HDPE or LLDPE), for example[7]. The mechanism of degradation of polymers has generally been described as free radical in the case of a thermal process without catalyst. However, when catalysts are used, it is generally ionic mechanism[7]. When catalysts are utilized in the pyrolysis occur two kinds of decomposition mechanisms simultaneously: thermal cracking, which in turn can follow different mechanisms (random chain scission, scission the end of the chain and / or elimination of side groups) and catalytic cracking 45


Almeida, D., & Marques, M. F. (carbenium ions adsorbed on the catalyst surface, beta scission and desorption). As a result, a wide variety of products is generated, which in turn will react with each other resulting in a countless number of possible reaction mechanisms[30]. For the pyrolysis of polyolefins, the degradation mechanism occurs by random chain scission, where free radicals are generated propagating chain reactions and thus resulting in the cracking of polymers in a wide range of hydrocarbons that make up liquid and gaseous fractions[32]. Several factors influence the process and the most important are: residence time, temperature and the type of pyrolysis agent. When the residence time and temperature increase, the composition of the obtained product shifts to more thermodynamically stable compounds[2,8,20,32]. The pyrolysis products can be used as an alternative fuel or as a source of chemicals[30]. The composition of the product also depends on the presence of catalysts (including concentrations and types). Higher temperatures decrease the yield of hydrogen, methane, acetylene and aromatic compounds, whereas lower temperatures favor the generation of gas products[32]. Previous experiments to evaluate the polymer degradation process are important because they provide information on the feasibility of recycling these polymers raw materials and even fuels. However, most studies are focused on pyrolysis of pure polymers and unmixed[7]. 1.1.1 Thermal pyrolysis The pyrolysis of waste plastics involves the thermal decomposition in the absence of oxygen / air. During the pyrolysis, the polymer materials are heated to high temperatures and thus, their macromolecules are broken into smaller molecules, resulting in the formation of a wide range hydrocarbons. The products obtained from the pyrolysis can be divided into non-condensable gas fraction, liquid fraction (consisting of paraffins, olefins, naphthenes and aromatics) and solid waste. From the liquid fraction can be recovered hydrocarbons in the gasoline range (C4-C12), diesel (C12-C23), kerosene (C10-C18) and motor oil (C23-C40)[1,3,18,20,33-35]. The thermal cracking usually produces a mixture of low value hydrocarbons having a wide variety of products, including hydrogen to coke. In general, when the pyrolysis temperature is high, there is increased production of non‑condensable gaseous fraction and a lower liquid fuels such as diesel. The yield and composition of the products obtained are not controlled only by the temperature but also the duration of the reaction[33]. The thermal pyrolysis proceeds according to the radical chain reactions with hydrogen transfer steps and the gradual breakdown of the main chain. The mechanism involves the stages of initiation, propagation and / or free radical transfer followed by β chain scission and termination[20,34,36]. This mechanism provides many oligomers by hydrogen transfer from the tertiary carbon atom along the polymer chain to the radical site[18]. The thermal cracking is more difficult for the high density polyethylene (HDPE), followed by the low density (LDPE) and then by polypropylene (PP)[20]. This is due to high content of tertiary carbons of PP. 46

The initiation step comprises homolytic breaking of carbon-carbon bond, either by random chain scission as by cleavage at the end of the chain, resulting in two radicals[36,37]. For PP and PE the chain scission occurs at random[37]. This step is followed by hydrogen transfer reactions intra / intermolecular forming more stable radicals secondary. These intermediate radicals can be submitted to break the carbon-carbon bond by scission β to produce compounds saturated or with unsaturated terminal and new radicals. The transfer of intra / intermolecular hydrogen depend on the experimental conditions, the first of which leads to an increase in the production of olefins and diolefins, paraffins results in the second[34,36,37]. The termination reactions can occur, for example, by disproportionation, which can produce different olefins and alkanes or a combination of radicals can lead to the same products. Branched products can be formed from the interaction between two secondary radicals or between a secondary radical with a primary[36,37]. As a consequence of these mechanisms, the thermal pyrolysis leads to a wide distribution of hydrocarbon, a C5‑C80 range, each fraction being mainly composed of diene, 1-olefin and n-paraffin. At high temperatures hydrogen is formed in significant amounts. Products obtained by thermal cracking are of limited commercial value, especially being applied as fuel. For heavy oils, it has been proposed its use as a wax[36]. Obtaining this wide range of products is one of the major drawbacks of this technique, which requires temperatures of 500 °C to 900 °C. These factors severely limit its applicability and increase the cost of recycling raw material of plastic waste[23]. 1.1.2 Catalytic pyrolysis The thermal pyrolysis requires high temperatures due to the low thermal conductivity of polymers[20], which is not very selective and a possible solution to reduce these reaction conditions is the use of catalyzed pyrolysis. Catalytic pyrolysis is an alternative to the recycling of pure or mixed plastics waste[30]. The catalyst can promote: • decomposition reactions at low temperatures with lower energy consumption[15,20,36,38,39]; • reduced costs[40]; • increase the yield of products with higher added value[20,38,40]; • increase the process selectivity[39,41]; • faster cracking reactions, leading to smaller residence times and reactors with smaller volumes[36]; • inhibiting the formation of undesirable products[36]; • inhibiting the formation of products consisting primarily of cyclic hydrocarbons, aromatic and branched, in the case of polyolefins catalytic cracking[36]; • obtain liquid products with a lower boiling point range[33].

Homogeneous and heterogeneous catalyst systems have been employed in the cracking polymers. In general, heterogeneous catalysts have been more used due to the ease of their separation and recovery of the reaction[36,39]. Polímeros, 26(1), 44-51, 2016


Thermal and catalytic pyrolysis of plastic waste The homogeneous catalysts especially used are Lewis acids, as AlCl3, fused metal tetracloroaluminatos (M (AlCl4) n), where the metal may be lithium, sodium, potassium, magnesium, calcium or barium and n can be 1 or 2)[36]. A wide variety of heterogeneous catalysts has been used and among them are: conventional solid acids (such as zeolites, silica-alumina, alumina and FCC catalysts (Fluid Catalytic Cracking)), mesostructured catalysts (such as MCM-41 etc.), nanocrystalline zeolites (such as n-HZSM-5), among others[25,35,36,39]. Many studies have been carried out describing the cracking of pure polyolefins over various solid acids such as zeolites, clays, among others. The use of zeolites has been shown to be effective in improving the quality of products obtained in the pyrolysis of polyethylene and other addition polymers. The acidity of their active sites and its crystalline microporous structure (textural properties) favor hydrogen transfer reactions and thereby make them suitable for obtaining high conversions of gas at relatively low temperatures, between 350 and 500 °C[11,18,22,41-44]. That is, these features allow milder operating conditions (lower temperatures and reaction times) than a thermal pyrolysis[4,25,30,45]. Differences in the catalytic activity of these solids are related to their acidic properties, especially the strength and number of acidic sites. The properties of these solid structures, as the specific area, particle size and pore size distribution, also have a crucial role in their performance, they control accessibility of voluminous molecules of the polyolefin internal catalytically active sites. While most work on catalytic cracking of polymers has been performed with pure polymers, it is accepted that the decomposition process can be affected by the presence of contaminants as well as chemical changes that occur in the polymer structure during use[11,20,21,34,42,46]. As mentioned, the catalyst pore size and acidity are important factors in the catalytic cracking of polymers[40,43,47]. Generally, the level of catalytic activity in the polyolefin pyrolysis increases with increasing the number of acidic sites. Thus, it is known that zeolite catalysts achieve higher conversions acids non-zeolitic catalysts[42]. The mechanism of this process which involves the formation of a carbenium ion (isomerization, random chain scission and β cleavage, hydrogen transfer, oligomerization / alkylation, aromatization) is influenced by the strength, density and distribution of the acid sites of the catalyst. This determines the products obtained in these reactions. Solid acid catalysts such as zeolites, favor hydrogen transfer reactions due to the presence of many acid sites[11,18,22,36,42,44]. The acid strength of the solid is characterized by the presence of Lewis or Brønsted acid sites. In the case of crystalline solid acids, it is believed that most of the acid sites are located inside the pores of the material, as in the case of zeolites[11,42]. Cracking is processed either by random chain scission (medium or weak acidity), for scission at the end of the chain (strong acidity) to give waxes and distillates (gasoil, gasoline) or light hydrocarbons (C3-C5 olefins), respectively. These primary cracking products may be removed from the reaction medium or subjected to secondary reactions (such as oligomerization, cyclization and aromatization). Polímeros, 26(1), 44-51, 2016

The relative extent of these reactions is connected to the acidity and properties of catalyst, but also to experimental variables employed (such as reactor type, temperature, residence time, etc.)[36]. Catalysts having acidic sites on the surface and with the possibility of donating hydrogen ion increase rate of the isomerization products and increase the yield of hydrocarbon isomers and the quality of the fuel formed. Catalysts containing strong acid sites, higher density, are more effective in cracking polyolefins. However, the strong acidity and high pore size cause rapid deactivation of the catalyst. Thus, according to literature, it is preferable to carry out the pyrolysis of polyolefins in the presence of a catalyst with light acidity and long life[33]. Other types of catalysts which may be used in the pyrolysis process are catalysts with Lewis acid sites which are electron pair acceptors. As examples of such catalysts, there are AlCl3, FeCl3, TiCl4 and TiCl3, which are strong Lewis acids[47]. These catalysts may be dissolved in molten polymer, which substantially increases the cracking efficiency while reducing its consumption. These types of catalysts have acidic sites on their surfaces that change the charge distribution in the carbon chain, making them capable of abstracting hydride ions of hydrocarbons to produce carbonium ions. This increases the catalytic effect, enabling a reduction in pyrolysis temperature and promoting the generation of ions for olefinic and aromatic compounds[32]. However, the cost of the catalyst can greatly affect the economy of the process, even if it shows a good performance. To reduce this cost and make it even more attractive process, you can reuse the catalyst or use it in smaller quantities[23,42,48]. The biggest problem in the use of catalysts in the pyrolysis of plastics is that coke formation deactivates the catalyst over time, thereby decreasing its life cycle[33].

1.2 Comparison between thermal and catalytic pyrolysis Seo et al.[49] studied the catalytic degradation of HDPE using a batch reactor at a temperature of 450 °C. As shown in Table 1, the pyrolysis performed with the zeolite ZSM-5 had higher yield of the gaseous fraction and smaller liquid fraction when compared with thermal cracking. According to the authors, this is explained by the properties of the catalyst. Most zeolites, including ZSM-5, showed excellent catalytic efficiency in cracking, isomerization and aromatization due to its strong acidic property and its microporous crystalline structure. The ZSM-5 zeolite has a three-dimensional pore channel structure with pore size of 5.4 × 5.6 Å which allows an increased cracking of larger molecules, beyond Table 1. Yield in thermal and catalytic pyrolysis of HDPE with ZSM-5[49]. Product Yield (% wt.) Gas Fraction Liquid Fraction

Solid Fraction

Total C6-C12 C13-C23 >C23

Thermal Pyrolysis 13.0 84.0 56.55 37.79 5.66 3.0

Catalytic Pyrolysis 63.5 35.0 99.92 0.08 0.0 1.5

47


Almeida, D., & Marques, M. F. the high Si / Al ratio which leads to an increase in thermal stability and acidity. Thus, initially degraded material on the external surface of the catalyst can be dispersed in the smaller internal cavities of the catalyst thus decomposed gaseous hydrocarbons (molecules with smaller sizes). Marcilla et al.[34] also used a batch reactor to evaluate the thermal and catalytic pyrolysis of HDPE and LDPE with HZSM-5 catalyst. The processing temperature was 550 °C and the results are shown in Table 2. As can be seen, the condensable products were the major fraction for the thermal process and no solid fraction (coke) was detected. For the catalytic process an increase of the gas fraction, and this is due to the HZSM-5 catalyst present, which has strong and weak acid sites and an average pore size small. As mentioned above, this facilitates cracking leading to compounds with small sizes (gas fraction). The results for the batch reactor are similar. However, there are studies where the values for each product obtained are different. This is because in this type of reactor the heat transfer is not as favored and, consequently, other factors such as the size and quantity of the sample or the carrier gas flow can determine the type of product formed. Moreover, in such reactors the extent of secondary reactions is smaller than the fluidized bed reactor. Using fixed beds where polymer and catalyst are contacted directly leads to problems of blockage and difficulty in obtaining intimate contact over the whole reactor. Without effective contact the formation of large amounts of residue are likely, and scale-up to industrial scale is not feasible[15]. The low thermal conductivity and high viscosity of the plastic may lead to a difficulty in mass transfer and heat. These factors influence the distribution of products, in conjuction with the operating conditions[50].

1.3 Zeolites Zeolites are microporous crystalline aluminosilicates of the elements of group 1A or 2A (especially sodium, potassium, magnesium and calcium), whose chemical composition can be represented as follows: M2 / nO.Al2O3.ySiO2.wH2O, where y varies from 2 to 10, n is the valence of the cation and w is the amount of structural water[36]. Currently it is known the existence of minerals which have all essential requirements to be classified as zeolites, however, instead of aluminum (Al) and silicon (Si) occupying the tetrahedral positions are present elements such as phosphorus (P), beryllium (Be), among others[51,52]. They are composed of tetrahedra of SiO4, AlO4 and PO4 as primary structural units, which are linked through oxygen atoms. Each oxygen atom is shared by two silicon or aluminum atoms, thus giving rise to a three-dimensional microporous structure[46,53]. The combination of these Table 2. Yield of the thermal and catalytic pyrolysis of LDPE and HDPE with HZSM-5[34]. Product Yield (% wt.) Gas Fraction Liquid Fraction/wax Solid Fraction

48

LDPE

HDPE

LDPEHZSM-5

HDPEHZSM-5

14.6 93.1

16.3 84.7

70.7 18.3

72.6 17.3

-

-

0.5

0.7

two primary structures is found in the common zeolites, developing cavities of various shapes and sizes which are interconnected[42,51,53,54]. The AlO4 tetrahedron has a negative charge of -1, because the aluminum has a valence of +3, which is less than the valence of +4 silicon. This charge is balanced by cations of alkali metals or alkaline earth metals (typically Na +, K +, Ca +2 or Mg +2) present inside the porous zeolite structure by means of cation exchange, may be replaced by other cations. When these cations are exchanged for protons, zeolite acid sites are formed. This exchange allows modification of the original properties of zeolites. The acidity of the zeolite can be the Brønsted acid type, proton donors or Lewis acid type, pair of electron acceptor[46,53]. These channels and cavities are occupied by ions, water molecules or other adsorbates which, due to high mobility, allow the ion exchange[51,53]. The pore size corresponding to two-dimensional opening zeolite is determined by the number of tetrahedral atoms connected in sequence. The three-dimensional interactions lead to the most different geometries, forming from large internal cavities to a series of channels crossing the whole zeolite[55]. The pores of zeolites function as molecular sieves, blocking the free diffusion of large, bulky molecules inside the internal surface of the catalyst[41,54]. These molecular sieves combine high acidity with selectivity form. That is, are selective to separate molecules according to their shape and / or size, besides having a high specific area and high thermal stability to catalyze a variety of hydrocarbon reactions, including the cracking of polyolefins. The reactivity and the selectivity of zeolites as catalysts are determined by its high number of active sites, which are caused by an imbalance of charge between the silicon and aluminum atoms in the crystal, making the zeolite of the structural unit has a charge balance total least one[42,51]. However, the process of rupture of the polymer molecules starts on the external surface of the zeolites, since the polymer chains must be broken before penetrating the internal pores of the zeolites, due to its small pore size. The zeolites have a specific pore size and the access of polymer molecules to internal reactive sites of the catalyst, as well as the final products within the pores are limited by their size. As mentioned, the catalyst pore size and acidity are important factors in the catalytic cracking of polymers[40,43,47]. Generally, the level of catalytic activity in the pyrolysis of polyolefins increases with increasing the number of acidic sites. Thus, it is known that zeolite catalysts achieve higher conversions than non-zeolitic catalysts acids[42]. In addition, branching of the polymer or end chain of polyethylene can penetrate the pores of the zeolites, reacting the acid sites located there and so increasing the activity[34]. During the catalyzed pyrolysis, the polymer melts and is dispersed around the catalyst. The molten polymer is drawn into the spaces between the particles and therefore the active sites on the external surface of the catalyst. Reactions at the surface produce a low molecular weight materials, which are sufficiently volatile at the temperature of the reaction can diffuse through the polymer film as a product or may react even more in the pores. These reactions proceed via carbocation as transition state. The reaction rate is governed both by the nature of the carbocation formed as the nature and strength of the acid sites involved in catalysis. Regardless Polímeros, 26(1), 44-51, 2016


Thermal and catalytic pyrolysis of plastic waste of how the carbocation is formed, it may be subjected to any of the following methods: load isomerization, the isomerization chain, hydride transfer, transfer of alkyl groups and formation and breaking of carbon-carbon bonds. As a result of this complex procedure, the product distribution reflects the action of the catalyst, which in turn is influenced by the size of its pores and for its chemical composition[34,56]. The catalytic decomposition of the polyethylene occurs at the carbenium ion mechanism. The initial step occurs either by abstraction of the hydride ion (for Lewis acid sites) or by addition of a proton (the Brønsted acid sites) in the C-C bonds of polyethylene molecules, or by thermal decomposition of polyolefins. Successive scission of the main chain occur to produce fragments having lower molecular weights than that of polyethylene. The resulting fragments are cracked or desidrociclizados in subsequent steps[18]. The acid sites on the catalyst surface are responsible for the initiation of the carbocationic mechanism, which induces the degradation of polyethylene and polypropylene. As mentioned above, these acid sites are originated the generated load imbalance when AlO4– is incorporated in the structure of zeolites. The content of AlO4– determines the number of acid sites in the catalyst while topological factors related to its crystalline or amorphous structure influence the strength of these acidic centers. Textural characteristics control the access of molecules that are reacting in the catalytic sites. This accessibility is important in catalyzed reactions involving large molecules such as polymers[21,57]. For presenting a microporous structure, zeolites have a higher internal surface than the external surface and this enables the mass transfer between these surfaces. However, the pore size is an important factor in this transfer, because only molecules with sizes smaller than these pores can enter or leave these spaces, which vary from one to another zeolite[53]. Some chemical and physical characteristics of zeolites ensures them their catalytic capacity. Among these characteristics can be cited: high specific area and adsorption capacity; active sites (which may be acidic) whose strength and concentration can be directed to a specific application; size channels and cavities compatible with the size of many molecules and a network of canals and cavities that provides you with a selectivity of shape, selectivity to the reactant, product and transition state species[53]. One of the factors that can affect the catalytic activity of zeolites is their deactivation by coke deposition on their channels. However, this coke formation rate depends on several factors, including: structure and acidity of the pores and the reaction conditions (such as temperature, pressure and nature of the reactants)[53]. The synthetic zeolites present some advantages and disadvantages in relation to natural zeolites. Among the advantages may be mentioned the purity, uniformity in size and shape of the channels and cavities, and a pre-defined chemical composition. The disadvantage has been their high cost and because of this, the synthetic zeolites are mainly intended for specific applications, where there is a need for a uniform composition and structure, for example, in the petroleum cracking process. Already the natural have a greater abundance and a lower cost of production, particularly if used in its in natura or if they require little beneficiation complex processes[51]. Polímeros, 26(1), 44-51, 2016

2. Conclusions Consumption of plastics has increased over the years and the concern with their waste generated too. Because of this many studies have been done with the aim to recover or recycle the waste. Pyrolysis has been effective compared to other disposal methods, because it can reuse the energy and the raw materials contained in those waste, reducing thereby the environmental impacts caused by the inadequate disposal of these waste plastics. The pyrolysis process may be thermal or catalytic. Thermal degradation occurs by radical mechanism, and as a result of this mechanism the products formed have a broad distribution of the number of carbon atoms in the main chain. In this type of the endothermic process due to the low thermal conductivity of polymers, there is a need for high temperatures. Because of that there is a high expenditure of energy. In order to decrease this temperature, catalysts may be used. With the catalytic pyrolysis, the products obtained have a more narrow distribution of the number of carbon atoms being directed to more specific products. The composition and amount of the obtained products are listed as type of catalyst used. Furthermore, the catalytic reaction decreases the degradation time and the fraction of solid waste formed. Generally, the catalysts used in the catalytic degradation are solid acids such as zeolites. This type of degradation involves production of the intermediate carbenium ion by hydrogen transfer reactions. Zeolites used favor these reactions due to their sites acids that help in the process of breaking the polymer macromolecules. This breaking process begins on the surface of the zeolite, because the polymer needs to be broken into smaller molecules before entering the internal pores of these solids, due to the small size of their pores. Zeolites have a specific molecular pore size and access of such molecules to catalytic reactive sites, as well as growth of the final products within such pores is limited by its size. The other experimental parameters such as temperature, reaction time, reactor type and flow of carrier gas also influence the composition of the products obtained. Pyrolysis can be carried out either for pure polymers or for polymer blends.

3. References 1. Mastral, J. F., Berrueco, C., & Ceamanos, J. (2007). Theoretical prediction of product distribution of the pyrolysis of high density polyethylene. Journal of Analytical and Applied Pyrolysis, 80(2), 427-438. http://dx.doi.org/10.1016/j.jaap.2006.07.009. 2. Abbas-Abadi, M. S., Haghighi, M. N., & Yeganeh, H. (2012). The effect of temperature, catalyst, different carrier gases and stirrer on the produced transportation hydrocarbons of LLDPE degradation in a stirred reactor. Journal of Analytical and Applied Pyrolysis, 95, 198-204. http://dx.doi.org/10.1016/j. jaap.2012.02.007. 3. Arabiourrutia, M., Elordi, G., Lopez, G., Borsella, E., Bilbao, J., & Olazar, M. (2012). Characterization of the waxes obtained by the pyrolysis of polyolefin plastics in a conical spouted bed reactor. Journal of Analytical and Applied Pyrolysis, 94, 230-237. http://dx.doi.org/10.1016/j.jaap.2011.12.012. 49


Almeida, D., & Marques, M. F. 4. Coelho, A., Costa, L., Marques, M. M., Fonseca, I. M., Lemos, M. A. N. D. A., & Lemos, F. (2012). The effect of ZSM-5 zeolite acidity on the catalytic degradation of highdensity polyethylene using simultaneous DSC/TG analysis. Applied Catalysis A: General, 413-414, 183-191. http://dx.doi. org/10.1016/j.apcata.2011.11.010. 5. Abbas-Abadi, M. S., Haghighi, M. N., & Yeganeh, H. (2013). Evaluation of pyrolysis products of virgin high density polyethylene degradation using different process parameters in a stirred reactor. Fuel Processing Technology, 109, 90-95. http://dx.doi.org/10.1016/j.fuproc.2012.09.042. 6. Stelmachowski, M. (2010). Thermal conversion of waste polyolefins to the mixture by hydrocarbons in the reactor with molten metal bed. Energy Conversion and Management, 51(10), 2016-2020. http://dx.doi.org/10.1016/j.enconman.2010.02.035. 7. Miskolczi, N., & Nagy, R. (2012). Hydrocarbons obtained by waste plastic pyrolysis: comparative analysis of decomposition described by different kinetic models. Fuel Processing Technology, 104, 96-104. http://dx.doi.org/10.1016/j.fuproc.2012.04.031. 8. Demirbas, A. (2004). Pyrolysis of municipal of plastic wastes for recovery of gasoline-range hydrocarbons. Journal of Analytical and Applied Pyrolysis, 72(1), 97-102. http://dx.doi. org/10.1016/j.jaap.2004.03.001. 9. Valle, M. L. M., Guimarães, M. J. O. C., & Sampaio, C. M. S. (2004). Degradação de poliolefinas utilizando catalisadores zeólitas. Polímeros: Ciência e Tecnologia, 1(14), 17-21. http:// dx.doi.org/10.1590/S0104-14282004000100009. 10. Shah, S. H., Khan, Z. M., Raja, I. A., Mahmood, Q., Bhatti, Z. A., Khan, J., Farooq, A., Rashid, N., & Wu, D. (2010). Low temperature conversion of plastic waste into light hydrocarbons. Journal of Hazardous Materials, 179(1-3), 15-20. http://dx.doi. org/10.1016/j.jhazmat.2010.01.134. PMid:20172649. 11. Panda, A. K., Singh, R. K., & Mishra, D. K. (2010). Thermolysis of waste plastics to liquid fuel. A suitable method for plastic waste management and manufacture of value added products: a world prospective. Renewable & Sustainable Energy Reviews, 14(1), 233-248. http://dx.doi.org/10.1016/j.rser.2009.07.005. 12. Al-Salem, S. M., Lettieri, P., & Baeyens, J. (2009). Recycling and recovery routes of plastic solid waste (PSW): a review. Waste Management (New York, N.Y.), 29(10), 2625-2643. http:// dx.doi.org/10.1016/j.wasman.2009.06.004. PMid:19577459. 13. Spinacé, M. A. S., & De Paoli, M. A. (2005). A tecnologia da reciclagem de polímeros. Quimica Nova, 98(1), 65-72. http:// dx.doi.org/10.1590/S0100-40422005000100014. 14. Singhabhandhu, A., & Tezuka, T. (2010). The waste-to-energy framework for integrated multi-waste utilization: waste cooking oil, waste lubricating oil, and waste plastics. Energy, 35(6), 2544-2551. http://dx.doi.org/10.1016/j.energy.2010.03.001. 15. Lin, Y.-H., & Yang, M.-H. (2008). Tertiary recycling of polyethylene waste by fluidized-bed reactions in the presence of various cracking catalysts. Journal of Analytical and Applied Pyrolysis, 83(1), 101-109. http://dx.doi.org/10.1016/j. jaap.2008.06.004. 16. Huang, W.-C., Huang, M.-S., Huang, C.-F., Chen, C.-C., & Ou, K.-L. (2010). Thermochemical conversion of polymer wastes into hydrocarbon fuels over various fluidizing cracking catalysts. Fuel, 89(9), 2305-2316. http://dx.doi.org/10.1016/j. fuel.2010.04.013. 17. López, A., De Marco, I., Caballero, B. M., Adrados, A., & Laresgoiti, M. F. (2011). Deactivation and regeneration of ZSM-5 zeolite in catalytic pyrolysis of plastic wastes. Waste Management (New York, N.Y.), 31(8), 1852-1858. http://dx.doi. org/10.1016/j.wasman.2011.04.004. PMid:21530221. 18. Park, D. W., Hwang, E. Y., Kim, J. R., Choi, J. K., Kim, Y. A., & Woo, H. C. (1999). Catalytic degradation of polyethylene over 50

solid acid catalysts. Polymer Degradation & Stability, 65(2), 193-198. http://dx.doi.org/10.1016/S0141-3910(99)00004-X. 19. Lin, Y.-H., & Yang, M.-H. (2005). Catalytic reactions of postconsumer polymer waste over fluidized cracking catalysts for producing hydrocarbons. Journal of Molecular Catalysis A Chemical, 231(1-2), 113-122. http://dx.doi.org/10.1016/j. molcata.2005.01.003. 20. Achilias, D. S., Roupakias, C., Megalokonomos, P., Lappas, A. A., & Antonakou, E. V. (2007). Chemical recycling of plastic wastes made from polyethylene (LDPE and HDPE) and polypropylene (PP). Journal of Hazardous Materials, 149(3), 536-542. http://dx.doi.org/10.1016/j.jhazmat.2007.06.076. PMid:17681427. 21. Aguado, J., Serrano, D. P., San Miguel, G., Escola, J. M., & Rodríguez, J. M. (2007). Catalytic activity of zeolitic and mesostructured catalysts in the cracking of pure and waste polyolefins. Journal of Analytical and Applied Pyrolysis, 78(1), 153-161. http://dx.doi.org/10.1016/j.jaap.2006.06.004. 22. Serrano, D. P., Aguado, J., Escola, J. M., & Rodríguez, J. M. (2005). Influence of nanocrystalline HZSM-5 external surface on the catalytic cracking of polyolefins. Journal of Analytical and Applied Pyrolysis, 74(1-2), 353-360. http:// dx.doi.org/10.1016/j.jaap.2004.11.037. 23. Lin, H.-T., Huang, M.-S., Luo, J.-W., Lin, L.-H., Lee, C.-M., & Ou, K.-L. (2010). Hydrocarbon fuels produced by catalytic pyrolysis of hospital plastic wastes in a fluidizing cracking process. Fuel Processing Technology, 91(11), 1355-1363. http://dx.doi.org/10.1016/j.fuproc.2010.03.016. 24. López, A., De Marco, I., Caballero, B. M., Laresgoiti, M. F., & Adrados, A. (2010). Pyrolysis of municipal wastes: influence of raw material composition. Waste Management (New York, N.Y.), 30(4), 620-627. http://dx.doi.org/10.1016/j. wasman.2009.10.014. PMid:19926462. 25. Sakata, Y., Uddin, M.A., & Muto, A. (1999). Degradation of polyethylene and polypropylene into fuel oil by using solid acid and non-acid catalysts. Journal of Analytical and Applied Pyrolysis, 51(1-2), 135-155. 26. Lee, S.Y., Yoon, J.H., Kim, J.R., & Park, D.W. (2001). Catalytic degradation of polystyrene over natural clinoptilolite zeolita. Polymer Degradation and Stability, 74(2), 297-305. 27. Hernández, M. R., Garcia, A. N., & Marcilla, A. (2005). Study of the gases obtained in thermal and catalytic flash pyrolysis of HDPE in a fluidized bed reactor. Journal of Analytical and Applied Pyrolysis, 73(2), 314-322. http://dx.doi.org/10.1016/j. jaap.2005.03.001. 28. Buekens, A. (2006). Introduction to feedstock recycling of plastics. In J. Scheirs, & W. Kaminsky (Orgs.), Feedstock recycling and pyrolysis of waste plastics (pp. 3-42). Hoboken: John Wiley & Sons. 29. Silvério, F. O., Barbosa, L. C. A., & Piló-Veloso, D. (2008). A pirólise como técnica analítica. Quimica Nova, 31(6), 15431552. http://dx.doi.org/10.1590/S0100-40422008000600045. 30. Lopez-Urionabarrenechea, A., De Marco, I., Caballero, B. M., Laresgoiti, M. F., & Adrados, A. (2012). Catalytic stepwise pyrolysis of packaging plastic waste. Journal of Analytical and Applied Pyrolysis, 96, 54-62. http://dx.doi.org/10.1016/j. jaap.2012.03.004. 31. Singh, S., Wu, C., & Williams, P. (2012). Pyrolysis of waste materials using TGA-MS and TGA-FTIR as complementary characterization techniques. Journal of Analytical and Applied Pyrolysis, 94, 99-107. http://dx.doi.org/10.1016/j. jaap.2011.11.011. 32. Donaj, P. J., Kaminsky, W., Buzeto, F., & Yang, W. (2012). Pyrolysis of polyolefins for increasing the yield of monomers’ recovery. Waste Management (New York, N.Y.), 32(5), 840-846. http://dx.doi.org/10.1016/j.wasman.2011.10.009. PMid:22093704. Polímeros, 26(1), 44-51, 2016


Thermal and catalytic pyrolysis of plastic waste 33. Scheirs, J. (2006). Overview of commercial pyrolysis processes for waste plastics. In J. Scheirs, & W. Kaminsky (Orgs.), Feedstock recycling and pyrolysis of waste plastics (pp. 383434). Hoboken: John Wiley & Sons. 34. Marcilla, A., Beltrán, M. I., & Navarro, R. (2009). Thermal and catalytic pyrolysis of polyethylene over HZSM5 and HUSY zeolites in a batch reactor under dynamic conditions. Applied Catalysis B: Environmental, 86(1-2), 78-86. http:// dx.doi.org/10.1016/j.apcatb.2008.07.026. 35. Lee, K.-H. (2012). Effects of the types of zeolites on catalytic upgrading of pyrolysis wax oil. Journal of Analytical and Applied Pyrolysis, 94, 209-214. http://dx.doi.org/10.1016/j. jaap.2011.12.015. 36. Aguado, J., Serrano, D. P., & Escola, J. M. (2006). Catalytic upgrading of plastic wastes. In J. Scheirs, & W. Kaminsky (Orgs.), Feedstock recycling and pyrolysis of waste plastics (pp. 73-110). Hoboken: John Wiley & Sons. 37. Lee, K.-H. (2006). Thermal and catalytic degradation of waste HDPE. In J. Scheirs, & W. Kaminsky (Orgs.), Feedstock recycling and pyrolysis of waste plastics (pp. 129-160). Hoboken: John Wiley & Sons. 38. Murata, K., Brebu, M., & Sakata, Y. (2010). The effect of silicaalumina catalysts on degradation of polyolefins by a continuous flow reactor. Journal of Analytical and Applied Pyrolysis, 89(1), 30-38. http://dx.doi.org/10.1016/j.jaap.2010.05.002. 39. Liu, W., Hu, C., Yang, Y., Tong, D., Li, G., & Zhu, L. (2010). Influence of ZSM-5 zeolite on the pyrolytic intermediates from the co-pyrolysis of pubescens and LDPE. Energy Conversion and Management, 51(5), 1025-1032. http://dx.doi.org/10.1016/j. enconman.2009.12.005. 40. White, R. L. (2006). Acid-catalyzed cracking of polyolefins: primary reaction mechanism. In J. Scheirs, & W. Kaminsky (Orgs.), Feedstock recycling and pyrolysis of waste plastics (pp. 45-72). Hoboken: John Wiley & Sons. 41. Mastral, J. F., Berrueco, C., Gea, M., & Ceamanos, J. (2006). Catalytic degradation of high density polyethylene over nanocrystalline HZSM-5 zeolite. Polymer Degradation & Stability, 91(12), 3330-3338. http://dx.doi.org/10.1016/j. polymdegradstab.2006.06.009. 42. Ofoma, I. (2006). Catalytic pyrolysis of polyolefins. Atlanta: Georgia Institute of Technology. 43. Li, X., Shen, B., Guo, Q., & Gao, J. (2007). Effects of large pore zeolite additions in the catalytic pyrolysis catalyst on the light olefins production. Catalysis Today, 125(3-4), 270-277. http://dx.doi.org/10.1016/j.cattod.2007.03.021. 44. Miskolczi, N., & Bartha, L. (2008). Investigation of hydrocarbon fractions form waste plastic recycling by FTIR, GC, EDXRFS and SEC techniques. Journal of Biochemical and Biophysical Methods, 70(6), 1247-1253. http://dx.doi.org/10.1016/j. jbbm.2007.05.005. PMid:17602751. 45. Elordi, G., Olazar, M., Aguado, R., Lopez, G., Arabiourrutia, M., & Bilbao, J. (2007). Catalytic pyrolysis of high density polyethylene in a conical spouted bed reactor. Journal of Analytical and Applied Pyrolysis, 79(1-2), 450-455. http:// dx.doi.org/10.1016/j.jaap.2006.11.010.

Polímeros, 26(1), 44-51, 2016

46. Manos, G. (2006). Catalytic degradation of plastic waste to fuel over microporus materials. In J. Scheirs, & W. Kaminsky (Orgs.), Feedstock recycling and pyrolysis of waste plastics (pp. 193-208). Hoboken: John Wiley & Sons. 47. Kaminsky, W., & Zorriqueta, I.-J. N. (2007). Catalytical and thermal pyrolysis of polyolefins. Journal of Analytical and Applied Pyrolysis, 79(1-2), 368-374. http://dx.doi.org/10.1016/j. jaap.2006.11.005. 48. López, A., De Marco, I., Caballero, B. M., Laresgoiti, M. F., Adrados, A., & Aranzabal, A. (2011). Catalytic pyrolysis of plastic wastes with two different types of catalysts: ZSM-5 zeolite and Red Mud. Applied Catalysis B: Environmental, 104(3-4), 211-219. http://dx.doi.org/10.1016/j.apcatb.2011.03.030. 49. Seo, Y.-H., Lee, K.-H., & Shin, D.-H. (2003). Investigation of catalytic degradation of high-density polyethylene by hydrocarbons group type analysis. Journal of Analytical and Applied Pyrolysis, 70(2), 383-398. http://dx.doi.org/10.1016/ S0165-2370(02)00186-9. 50. Lin, Y.-H. (2009). Production of valuable hydrocarbons by catalytic degradation of a mixture of post-consumer plastic waste in a fluidized-bed reactor. Polymer Degradation & Stability, 94(11), 1924-1931. http://dx.doi.org/10.1016/j. polymdegradstab.2009.08.004. 51. Monte, M. B. M., & Resende, N. G. A. M. (2005). Zeolitas naturais. In A. B. Luz, & F. A. F. Lins, Rocha e minerais industriais: usos e especificações (pp. 699-720). Rio de Janeiro: CETEM. 52. Letichevsky, S. (2008). Síntese e caracterização das zeolitas mordenita, ferrierita e ZSM-5 nanocristalinas (Tese de doutorado). Pontifícia Universidade Católica do Rio de Janeiro, Rio de Janeiro. 53. Tourinho, R.R.C. (2009). Estudo da acidez de zeolitas impregnadas com platina utilizando reações de troca H/D com aromáticos e correlações lineares de energia livre (Dissertação de mestrado). Universidade Federal do Rio de Janeiro, Rio de Janeiro. 54. Aksoy, Y. Y. (2010). Characterization of two zeolites for geotechnical and geoenvironmental applications. Applied Clay Science, 50(1), 130-136. http://dx.doi.org/10.1016/j. clay.2010.07.015. 55. Braga, A. A. C., & Morgon, N. H. (2007). Descrições estruturais cristalinas de zeolitos. Quimica Nova, 30(1), 178-188. http:// dx.doi.org/10.1590/S0100-40422007000100030. 56. Pinto, F., Costa, P., Gulyurtlu, I., & Cabrita, I. (1999). Pyrolysis of plastic wastes 2. Effect of catalyst on product yield. Journal of Analytical and Applied Pyrolysis, 51, 57-71. 57. Hwang, E.-Y., Kim, J.-R., Choi, J.-K.; Woo, H.-C., & Park, D.-W. (2002). Performance of acid treated natural zeolitas in catalytic degradation of polypropylene. Journal of Analytical and Applied Pyrolysis, 62(2), 351-364. Received: Mar. 31, 2015 Revised: July 08, 2015 Accepted: Aug. 31, 2015

51


http://dx.doi.org/10.1590/0104-1428.1748

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

Mechanical and thermomechanical properties of polyamide 6/Brazilian organoclay nanocomposites Renê Anisio da Paz1, Amanda Melissa Damião Leite1*, Edcleide Maria Araújo1, Vanessa da Nóbrega Medeiros1, Tomás Jeferson Alves de Melo1 and Luiz Antônio Pessan2 Materials Engineering Department, Universidade Federal de Campina Grande – UFCG, Campina Grande, PB, Brazil 2 Materials Engineering Department, Universidade Federal de São Carlos – UFSCar, São Carlos, SP, Brazil

1

*amandamelissa.lins@yahoo.com.br

Abstract Polymer/clay nanocomposites are a new class of composites with polymer matrices where the disperse phase is a silicate with elementary particles that have at least one of dimensions in nanometer order. Polyamide 6/Brazilian organoclay nanocomposites were prepared by melt intercalation, and the mechanical, thermal and thermomechanical properties were studied. The structure and morphology of the nanocomposites were evaluated by X-ray diffraction (XRD) and transmission electron microscopy (TEM). It was verified by XRD and TEM analysis that all systems presented exfoliated structure predominantly. By thermogravimetry (TG), nanocomposites showed higher stabilities in relation to pure polymer. It was observed that the nanocomposites showed better mechanical properties compared to the properties of polyamide 6. The heat deflection temperature (HDT) values of the nanocomposites showed a significant increase in relation to pure polymer. Keywords: nanocomposites, polyamide 6, HDT, organoclays.

1. Introduction Polymer nanocomposites with clay are a new class of composites where the polymer matrix phase is dispersed silicate consisting of elementary particles which have at least one of its dimensions in the order of nanometers. The mineral particles commonly used in these materials are the smectite clays (montmorillonite, hectorite and saponite) in its particles lamellar morphology, with sides on the order of a micrometer and a thickness of approximately one nanometer[1,2]. From the 60s, the literature began to report the development of the first polymer/clay nanocomposites. From then until the present day, much attention has been given to polymer nanocomposites, especially those developed with layered silicates, due to the great need of modern materials of engineering and the fact that the pure polymers do not present the behavior or the properties required for certain application[3-15]. Organic/inorganic hybrids exhibit improved properties compared to the pure polymers or conventional composites, such as higher elastic modulus and tensile strength, higher resistance to solvents and flame resistance and good optical, magnetic and electric properties[4]. The improvement in the properties of these materials is achieved with a small load volume fraction (1-10%), and due to very high aspect ratio of the load, i.e., length/diameter ratio that is high and increases the interaction with the polymer. Moreover, the polymer nanocomposites have the additional advantage that they can be processed with techniques and equipment used for conventional polymers[5-15]. Polymer nanocomposites based on organoclays as a filler offer improved performance and opportunities for commercial

52

applications[16,17]. The key to significant enhancement in properties is to exfoliate the individual organoclay platelets into the polymer matrix to utilize their high aspect ratio and modulus[18]. The affinity between polymer matrix and organoclay is one of the most important factors in achieving good exfoliation; to a certain extent the affinity can be enhanced by optimizing the structure of the organoclay for a given polymer matrix. Previous studies have shown that semi-crystalline polyamides like nylon 6, nylon 66, nylon 11, nylon 12, etc. give rather good exfoliation[19,20]. Polyamide 6 (PA6) is one of the most used types of aliphatic polyamide. The main applications of PA6 are in fibres, films, and as injection-moulded engineering plastic. PA6 crystallizes fast, usually up to percentages in the range of 30-40%, providing a high modulus to the material even above the glass transition temperature (Tg). One property common to all polyamides is that they absorb water from the environment, both from the air and from liquid water[21]. The aim of this study was to prepare nanocomposites of polyamide 6 with three viscosity indexes and a Brazilian organoclay to evaluate the mechanical, thermal and thermomechanical properties. The use of Brazilian clay is the highlight and the differential of this work, whereas the literature as a whole frequently use commercial clay. The reason for the development of this work is the fact that the deposits of bentonite clay is abundant in South America. It is found in Brazil and therefore supplies the whole country with bentonite, and moreover it is essential to observe that the use of nanofillers brings benefits for the country and has a high technological and market importance.

Polímeros, 26(1), 52-60, 2016


Mechanical and thermomechanical properties of polyamide 6/Brazilian organoclay nanocomposites

2. Experimental It was used three polyamide 6 (PA6), Technyl C216 from Rhodia/SP, with viscosity index (VI)=134mL/g; and polyamide 6 from Polyform B300, with VI=140-160 mL/g and B400 with VI=235-265 mL/g, all in the form of granules of white coloration. The bentonite clay was Brasgel PA (sodium), CEC = 90 meq/100 g, provided by Bentonit União Nordeste (BUN), located in Campina Grande/PB/Brazil, in the form of powder passed in an ABNT 200 mesh sieve (D = 0.074 mm). The quaternary ammonium salt was Cetremide (hexadecyltrimethyl ammonium bromide). All materials containing polyamide were dried under vacuum at 80 °C for 24 hours. The organoclay was produced from cation exchange reaction, where the sodium ions present in the clay are exchanged for ammonium ions of the quaternary salt. The Na‑MMT as mixed in distilled water, was heated at 80 ± 5 °C and they were kept for 20 minutes with stirring to form a uniformly dispersed suspension. The salt equivalent to 1:1 CEC of Na-MMT as added into the dispersion. The mixture was stirred for more 20 minutes. After 24 h the mixture of bentonite and the salt were washed with distilled water for several times to remove the salt excess and it was dried at 60°C for 48 h, and finally, passed in a sieve 200 mesh detailing of the procedure is described by Díaz[22,23] and others authors[8,11,15,24]. In the nanocomposite preparation, before any processing step, all the materials with PA6 were dried in an oven with circulating air at 80°C for 1 h. Afterwards, these materials were kept in an oven under vacuum at 80°C for 24 h. The nanocomposites were prepared by two steps. Firstly, in order to assure a better dispersion of the fine clay powder in polyamide, a 1:1 PA6/organoclay master was previously produced in a Torque Rheometer Haake with internal mixer, at 240 °C and 60 rpm for 10 minutes. After, PA6/organoclay nanocomposites, containing 3 and 5 wt. (%) of clay, were melted in a corrotating twin-screw extruder operating at 240 °C and 250 rpm. After all the material has been extruded, granulated and dried at 80 °C in a vacuum oven for 24 hours, it was submitted to the process of injection molding in Allrounder injection Arburg 270/30 ton, with the following injection conditions: injection pressure of 38 MPa; temperature profile of 250 °C; mould temperature of 65 °C;

mould cooling of 20s; holding pressure of 32 MPa and injection speed of 27 cm3/s. The preparation of nanocomposites was according to literature[8,11,15,24]. The thermogravimetry analysis (TG) was performed on a Thermal Analyzer TGA Q500 (TA Instruments), using about 5 mg of sample heating rate of 20 °C min-1 and a sample holder of alumina. Samples were heated from room temperature to 900°C under nitrogen (N2) gas at a flow rate of about 50 ml min-1. Clays and nanocomposites were characterized by X-ray diffraction (XRD) in XRD‑6000 Shimadzu machine, using Kα radiation of the copper (λ = 1.542 Å), voltage 40 kV, 30 mA, scan 2θ between 2-30° and speed scanning at 2°C/min. The morphology of the nanocomposites were evaluated using the transmission electron microscope of the Philips CM120, operating at an acceleration voltage of 120 kV. Tensile tests were performed in accordance with ASTM D638-99 in a mechanical testing universal machine Instron 5569, with the deformation rate of 50 mm/min and the samples were conditioned into a desiccator for 48 hours before to testing. The properties determined were: yield stress, elongation at break and Young modulus. The results were obtained from ten (10) samples. The analysis of differential scanning calorimetry (DSC) have been made in the DSC Q20 TA Instruments. The samples were first heated from room temperature up to 260 °C and kept for 1 minute before cooling down to room temperature. A second heating was used to observe the melting behavior of the PA6. All heating and cooling steps were done at 10 °C/min. The Heat Deflection Temperature (HDT) was obtained according to ASTM D 648-01, HDT VICAT 6 P/N 6921, CEAST model, 1.82 MPa, at 120 °C/h, room temperature and up to 300 °C. The results were obtained from six (6) samples.

3. Results and Discussion Generally, the processing temperature of the polymeric materials is higher than 150 °C, near the thermal limit of the organic salts. The structure of the quaternary ammonium ions is usually used with the objective to improve the compatibility with specific polymer. However, this molecular structure also determines its thermal stability. Figure 1a and 1b illustrate the TG and DTG curves for untreated clays (MMT) and organoclay (OMMT). It is observed in Figure 1a that the

Figure 1. TG and DTG curves of (a) untreated clays and (b) Organoclay. Polímeros, 26(1), 52-60, 2016

53


Paz, R. A., Leite, A. M. D., Araújo, E. M., Medeiros, V. N., Melo, T. J. A., & Pessan, L. A. MMT clay presents weight loss that occurs in the range of 30 ºC to 160 ºC, corresponding to the loss of adsorbed water, and another the range of 400 ºC to 600 ºC corresponding to dehydroxylation the clay mineral with approximately 12% of weight loss. For organoclay (Figure 1b) the weight loss was ~ 26%, with a small weight loss below 100 °C corresponding to a residue of adsorbed water. It indicates that a small amount of clay was not modified by salt. A weight loss in the range of 180 ºC to 460 ºC with a maximum at 260 °C, corresponds to the decomposition of the quaternary ammonium salt and other weight loss in the range of 600 ºC to 750 ºC, corresponding to dehydroxylation of clay mineral.

water and gaseous products under 180 °C; 2) evaporation of organic substances between 200 °C and 500 °C; 3) aluminum‑silicate dehydroxylation between 500 °C and 700 °C and 4) evaporation of the organic residues between 700 °C and 1000 °C[25,26].

Through untreated and organoclay clays thermogravimetry analysis, the levels of water and organic salt incorporated were calculated. The untreated clays (MMT) presented a water content of approximately 10%. On the other hand, organoclay (OMMT) water content was around 1.2% and 22.59% of organic salt. These results indicated that the hydrophilic characteristic of the clay was reduced.

The results of TG and DTG of PA6 with three viscosity indexes and their nanocomposites with 3 and 5% clay are presented in Figures 2, 3 and Table 1. Three main characteristics are important: first, all samples exhibit the decomposition process, with the B300 5% system exhibiting the highest thermal stability; the second, the process of decomposition of all samples occurs in the same temperature range, except for the C216 3% sample that it starts to decompose before. This result implies that the decomposition of PA6 was little affected by the addition of clay in this rate; and finally, when the temperature was above 350 °C, the nanocomposites started to decompose at higher temperature than pure PA6, indicating increased thermal stability after incorporation of organoclay.

The thermal decomposition of organoclays in general can be divided into four regions: 1) evaporation of adsorbed

Similar decomposition behavior is exhibited by nanocomposites and can be attributed to the dispersion

Figure 2. TG curves of (a) polyamide 6 with three viscosity indexes and (b) their nanocomposites with 3 and 5% clay.

Figure 3. DTG curves of (a) polyamide 6 with three viscosity indexes and (b) their nanocomposites with 3 and 5% clay. 54

Polímeros, 26(1), 52-60, 2016


Mechanical and thermomechanical properties of polyamide 6/Brazilian organoclay nanocomposites of the clay layers within the matrix of PA6. The thermal stability enhancement of PA 6 after adding 5 wt% of clay was a consequence of the nano-scaled dispersion of layered clay, which resulted in oxygen and heat permeability reductions in the PA 6 matrix during the heating scans. The results of DTG are shown in the Figure 3a, b. It is observed that the TG shows a single peak with maximum weight loss, Tmax = 430°C and Tonset = 394°C for pure PA 6. The presence of organoclay in the PA6 matrix changed to Tmax ~ 480 °C e Tonset = 406°C in all nanocomposites, i.e., higher stabilities compared to pure polymer. This evidences that different clay contents influence the thermal stability of materials. This behavior is in agreement with that reported by literature[27] where they obtained similar results using systems with Cloisite 30B clay. Figure 4a shows the XRD patterns of the clay untreated clays (MMT) and of the organoclay (OMMT). Making a comparison of the diffractograms, it is possible to visualize the efficiency of organophilization by increasing the basal interplanar distance (d001) of the treated clay in relation to the untreated clay. The diffractograms of the untreated clay (MMT) shows characteristic peaks of bentonite containing accessory materials, such as quartz (Q) in the range of 22-30°, which occurs for all samples, and also a band at approximately 7.0°

that indicates the basal interplanar distance d001 of 12.63 Å, that is characteristic of montmorillonites containing Na+ ions in the structure. Analyzing the diffractograms of the treated clay with Cetremide salt (OMMT), it is seen the peak shifts to smaller angles and consequent expansion of layers to 20.83 Å, due to penetration of the carbon chains attached to the quaternary ammonium salt, which favors the electrostatic interactions with the matrix, and it facilitates the incorporation of the polymer, according to the literature[4,8,11,15]. Figure 4b illustrates XRD patterns of polyamide 6 (PA6) and systems of PA6/organoclay (OMMT). It is noticed that the peak of the organoclay with d001 = 20.83 Å disappeared when it was incorporated at 3 and 5% concentration by weight in the matrix of polyamide 6. These results can indicate that all systems exhibit exfoliated structure and/or partially exfoliated one[8-11,15]. It is interesting to note that the curves of the nanocomposites showed the same behavior, independently of the difference in molecular weight of PA. These results were confirmed by other works developed[8,10,11,15]. It can also be observed that in the range from 17° to 26°, there are some peaks for the crystalline forms (α- and γ-phases) of the polyamide 6, wherein the monoclinic phase has distinct reflections (α1 and α2). For the injected material, the reflection peak of γ -phase was more intense, as observed

Table 1. TG results obtained for polyamide 6 with three viscosity indexes and their nanocomposites with 3 and 5% clay. Sample

PA6C216

PA6B300

PA6B400

Clay content

Humidity

Organic Material

%

%

%

0 3 5 0 3 5 0 3 5

3.7 3.7 3.9 3.7 3.8 3.8 4.0 4.3 3.9

95.3 94.6 93.3 94.9 93.9 91.8 95.5 92.4 91.3

Decomposition Temperature (peak) (°C) 461.9 471.4 475.9 477.2 478.4 475.9 478.4 477.2 475.9

Residues % 0.9 1.7 2.8 1.4 2.3 4.4 0.5 3.3 4.8

Figure 4. XRD patterns of pure PA6 and nanocomposites: (a) clays, (b) polyamides 6 with 3 and 5% organoclay. Polímeros, 26(1), 52-60, 2016

55


Paz, R. A., Leite, A. M. D., Araújo, E. M., Medeiros, V. N., Melo, T. J. A., & Pessan, L. A. in the literature. The formation of these two phases depends mainly on the crystallization conditions or addition of specific loads. The form of the α structure is known as the most thermodynamically stable[8,10,15]. Photomicrographs of PA6 C216 nanocomposite with 3% of OMMT (Figure 5a) and B300 system with 5% clay (Figure 6b) show exfoliated morphologies with clay lamellae well distributed in the polymer matrix. The system PA6 C216 with 5% clay (Figure 5b) shows a morphology partially exfoliated with exfoliated areas and intercalated structures in small areas. PA6 B300 nanocomposite with OMMT 3% (Figure 7a) presents a morphology composed of some exfoliated lamellae and some agglomerates of clay dispersed

in the matrix. Already PA6 B400 nanocomposite with OMMT, 3% and 5% (Figure 7a and 7b) exhibit morphologies with exfoliated clay lamellae well distributed in the polymer matrix. According Fornes et al.[10], the high molecular weight polymer and therefore its higher viscosity contributes to a higher shear or energy for separating the lamellae of the clay. Already the matrixes of smaller viscosities indexes do not allow complete exfoliation. These results confirm the results of X-ray diffraction, that is, confirms that the use of the two techniques is important in interpreting the type of nanocomposites formed. The Table 2 presents HDT data obtained for the systems. The incorporation of clay into polymers generally

Figure 5. TEM photomicrographs of PA6 C216 nanocomposites with (a) 3%, and (b) 5% organoclay.

Figure 6. TEM photomicrographs of PA6 B300 nanocomposites with (a) 3%, and (b) 5% organoclay. 56

Polímeros, 26(1), 52-60, 2016


Mechanical and thermomechanical properties of polyamide 6/Brazilian organoclay nanocomposites

Figure 7. TEM photomicrographs of PA6 B400 nanocomposite with (a) 3%, and (b) 5% organoclay.

Table 2. HDT of polyamide 6 with three viscosities indexes and their nanocomposites with 3 and 5% clay. Samples C216

B300

B400

Clay content (%) 0 3 5 0 3 5 0 3 5

HDT (°C) 102.1 ± 8.4 152.6 ± 3.0 193.7 ± 0.8 89.4 ± 3.2 106.7 ± 1.8 129.4 ± 6.9 97.6 ± 0.8 110.0 ± 3.7 199.3 ± 1.1

increases the material stiffness and consequently increases the HDT values. In this case, clay considerably increased the HDT values of nanocomposites mainly with the 5% clay. In general, the HDT of polyamide was in the range of 96 °C and the average nanocomposite with 3% clay, in the range of 123 °C and 5% clay around 174 °C. From the point of view of application, the increase in this property is important for automotive, aeronautical industries, etc. The increases in systems B400 and C216 with 5% loading demonstrate the nucleating effect of the clay in the polymer. The heat deflection temperature (HDT) is the temperature at which a polymer sample deforms under a specific load. This property is applied in many aspects of engineering and manufacturing of products. The incorporation of nanoscale clay generally increases the HDT of PA6[11,12]. Shen et al. [28] studied the effect of organoclay in the mechanical and thermal properties of the nanocomposites of polyamide 6 reinforced with glass fiber, and observed that clay caused an increase of approximately 65% in the HDT PA6/glass fiber/clay system. The Figure 8 shows the DSC curves of polyamide 6 and its nanocomposites. In Figure 8a it can be observed that the Polímeros, 26(1), 52-60, 2016

B300, B400, B300 3%, C216 3% and B300 5% compositions have different crystalline behaviors of other compositions. In Figure 8b, the crystallization temperatures practically unchanged for all samples. Already in Figure 8c, it can be observed that the melt peak shows a bimodal pattern of the endothermic peak in relation to pure PA6, and this pattern is seen with more intensity for compositions with increasing clay content. The endothermic peak Tm1 refers to α crystalline form of PA6. The presence of a second endothermic peak Tm2, with temperatures slightly below the Tm1, is related to the γ crystalline form of PA6. These melt temperatures (Tm1 and Tm2) suggest that α and γ crystalline forms coexist in the nanocomposite. The γ crystalline form may be associated with a decrease in the degree of packing of the crystalline domains. Therefore, the lower Tm values of the nanocomposites can be attributed changes in lamellar size and distribution of the crystallites of PA6[29-31] Table 3 shows that the presence of the clay had minor changes in the crystallization temperature (Tc) when compared to Tc of the pure polyamide. Probably the clay did not act as nucleating agent which can also be evidenced by the onset temperature of crystallization of the nanocomposites in relation to pure polymer. In general, it can be observed that the crystallinity degree (Xc) of the nanocomposites decreased in relation to pure polymer. This can be attributed to interaction of the clay with the phase of the polyamide that restricted the mobility of the polymer chains hindering the crystallization process[33]. Table 4 shows the results of mechanical properties of tensile polyamide 6 with three viscosities indexes and their nanocomposites with 3 and 5% of organoclay. It can be observed that the nanocomposites showed better mechanical properties under tensile as compared properties of polyamide 6, i.e., the clay probably acted as reinforcing filler increasing the rigidity of the system, as it can be verified by the elastic modulus and yield stress. According 57


Paz, R. A., Leite, A. M. D., Araújo, E. M., Medeiros, V. N., Melo, T. J. A., & Pessan, L. A.

Figure 8. DSC curves (a) first heating, (b) cooling and (c) second heating of polyamide 6 with three viscosities indexes and their nanocomposites with 3 and 5% clay.

Table 3. Melting and crystallization parameters obtained for polyamide 6 with three viscosities indexes and their nanocomposites with 3 and 5% organoclay. SAMPLE

PA6C216

PA6B300

PA6B400

Clay content (%) 0 3 5 0 3 5 0 3 5

Tm1

First Heating Xc1

ΔHm1

Tc

Cooling ΔHc1

Tonset

Tm2

(ºC) 224.3 221.4 222.9 222.9 219.8 222.9 225.5 222.5 222.6

(%) 46.5 37.2 33.6 34.3 42.1 35.8 47.3 32.6 30.9

(J/g) 87.5 69.9 66.4 64.4 61.6 70.9 88.9 63.2 61.2

(ºC) 190.8 190.7 188.6 191.0 189.5 188.3 190.3 187.3 186.2

(J/g) 66.8 65.2 59.2 65.7 59.8 58.3 64.9 58.3 57.6

(ºC) 195.9 195.4 194.1 196.4 194.3 193.3 196.1 191.7 191.4

(ºC) 222.1 223.0 221.3 221.6 221.6 221.4 221.9 220.8 221.1

Second Heating Xc2 ΔHm2 (%) 38.2 36.1 32.6 32.1 36.2 33.5 39.7 32.7 33.7

(J/g) 71.9 67.8 64.6 60.4 70.1 66.3 74.7 63.4 66.8

Tm = Melting temperature taken at the melt peak; ∆Hc = Crystallization Enthalpy; ΔHm = Heat of fusion, calculated from the melting peak; Tonset = onset temperature. Xc = Degree of crystallinity obtained by DSC, taken from ∆HF/∆HF100%; ∆HF100 = heat of fusion for PA6, 100% crystalline, 188 J/g[32].

to Sinha Ray and Okamoto[9] and Alexandre and Dubois[13] and, the system formed by polymer/clay containing low clay content (<10%) exhibit better mechanical properties compared to pure polyamide. This increase is because of the stronger interfacial interaction between the matrix and 58

the clay layers, especially when they are more dispersed 10. On the other hand, the elongation at break decreased, as expected, due to the particles of the clay acted as a barrier preventing the mobility of the polymer chains and reducing the elongation[15,24]. Polímeros, 26(1), 52-60, 2016


Mechanical and thermomechanical properties of polyamide 6/Brazilian organoclay nanocomposites Table 4. Mechanical properties obtained of polyamide 6 with three viscosities indexes and their nanocomposites with 3 and 5% organoclay. Samples

C216

B300

B400

Clay Content

Module

(%)

(GPa)

0

at break

3.3±0.0

67.7±1.1

(%) 38.5±1.7

3

3.5±0.1

66.4±1.5

25.6±2.0

5 0

4.8±0.1 3.6 ±0.0

81.2±4.7 75.1±1.6

8.1±0.8 12.1±0.5

3

3.9 ±0.1

86.0±1.2

7.7±0.3

5 0

4.1 ±0.2 3.7±0.1

85.1±1.1 72.4±0.5

4.6±0.4 14.2±0.9

3

3.8±0.2

75.9±1.1

10.3±0.6

5

4.0±0.2

82.3±1.4

9.9±0.5

4. Conclusions Nanocomposites of polyamide 6 with three viscosities indexes and a Brazilian organoclay were produced and the mechanical, thermal and thermomechanical properties were evaluated. By XRD, the peak of the organoclay disappeared when incorporated into polyamide 6 indicating that all systems had exfoliated and/or partially exfoliated structures. TEM photomicrographs confirmed these structures. By TG, nanocomposites showed higher thermal stability in relation to pure polymer. The nanocomposites showed HDT values significantly higher than that of pure polyamide 6. The nanocomposites with 5% clay had better tensile mechanical properties when compared to the nanocomposites with 3% and pure polyamides. In other words, the clay acted as reinforcing filler increasing the rigidity of the system. The mechanical properties (modulus and yield stress) increased with the presence of organoclay. The three viscosities indexes of PA6 and two contents of clay influenced the structure and properties of the systems, ie, the Brazilian clay had an important role in the final properties of HDT because there was an 100% increase in the system HDT in comparison with the pure polymer.

5. Acknowledgements The authors thank Rhodia/SP, Bentonit União Nordeste (BUN), DEMa/UFCG and DEMa/UFSCar, CAPES/PNPD, MCTI/CNPq for the financial support.

6. References 1. Shi, H., Lan, T., & Pinnavaia, T. J. (1996). Interfacial effects on the reinforcement properties of polymer organoclay nanocomposite. Chemistry of Materials, 8(8), 1584-1587. http://dx.doi.org/10.1021/cm960227m. 2. Souza Santos, P. (1989). Ciência e tecnologia de argilas. 2nd ed. São Paulo: Edgar Blucher. 3. Fornes, T. D., & Paul, D. R. (2003). Crystallization behavior of nylon 6 nanocomposites. Polymer, 44(14), 3945-3961. http:// dx.doi.org/10.1016/S0032-3861(03)00344-6. 4. García-López, D., Fernández, J. F., Merino, J. C., & Pastor, J. M. (2013). Influence of organic modifier characteristic on the mechanical properties of polyamide 6/organosepiolite nanocomposites. Composites. Part B, Engineering, 45(1), 459-464. http://dx.doi.org/10.1016/j.compositesb.2012.09.087. Polímeros, 26(1), 52-60, 2016

Elongation Yield stress (MPa)

5. Pavlidou, S., & Papaspyrides, C. D. (2008). A review on polymer–layered silicate nanocomposites. Progress in Polymer Science, 33(12), 1119-1198. http://dx.doi.org/10.1016/j. progpolymsci.2008.07.008. 6. Mészáros, L., Deák, T., Balogh, G., Czvikovszky, T., & Czigány, T. (2013). Preparation and mechanical properties of injection moulded polyamide 6 matrix hybrid nanocomposite. Composites Science and Technology, 75(11), 22-27. http:// dx.doi.org/10.1016/j.compscitech.2012.11.013. 7. Barbosa, R., Morais, D. D. S., Nóbrega, K. C., Araújo, E. M., & Mélo, T. J. A. (2012). Influence of processing variables on the mechanical behavior of HDPE/clay nanocomposites. Materials Research, 15(3), 477-482. http://dx.doi.org/10.1590/ S1516-14392012005000054. 8. Paz, R. A., Araújo, E. M., Pessan, L. A., Melo, T. J. A., Leite, A. M. D. & Medeiros, V. N. (2012). Influence of molecular weight of polyamide 6 in obtaining of nanocomposites with national organoclay. Materials Science Forum, 727, 1711-1777. Retrieved in 15 April 2014, from 10.4028/www.scientific.net/ MSF.727-728.1711 9. Sinha Ray, S., & Okamoto, M. (2003). Polymer/layered silicate nanocomposites: a review from preparation to processing. Progress in Polymer Science, 28(11), 1539-1568. http://dx.doi. org/10.1016/j.progpolymsci.2003.08.002. 10. Fornes, T. D., Yoon, P. J., Keskkula, H., & Paul, D. R. (2001). Nylon 6 nanocomposites: the effect of matrix molecular weight. Polymer, 42(25), 9929-9940. http://dx.doi.org/10.1016/S00323861(01)00552-3. 11. Paz, R. A., Leite, A. M. D., Araújo, E. M., Melo, T. J. A., & Pessan, L. A. (2010). Avaliação do comportamento térmico por DSC na Região da pele e do Núcleo de Amostras Injetadas de Nanocompósitos de Poliamida 6/Argila Organofílica. Polímeros, 20(4), 258-262. http://dx.doi.org/10.1590/S010414282010005000043. 12. Kojima, Y., Usuki, A., Kawasumi, M., Okada, A., Fukushima, Y., Kurauchi, T., & Kamigaito, O. (1993). Mechanical properties of polyamide 6-clay hybrid. Journal of Materials Research, 8(5), 5-11. http://dx.doi.org/10.1557/JMR.1993.1185. 13. Alexandre, M. & Dubois, P. (2000). Polymer-layered silicate nanocomposites; preparation, properties and uses of a new class of montmorillonite. Materials Science and Engineering, 28, 1-63. http://dx.doi.org/S0927796X(00)000127. 14. Ayres, E., & Oréfice, R. L. (2007). Nanocompósitos derivados de dispersões aquosas de poliuretano e argila: influência da argila na morfologia e propriedades mecânicas. Polímeros, 17(4), 339-344. http://dx.doi.org/10.1590/S0104-14282007000400015. 15. Paz, R. A., Araújo, E. M., Pessan, L. A., Melo, T. J. A., Leite, A. D., & Medeiros, V. N. (2012). Evaluation of impact strength 59


Paz, R. A., Leite, A. M. D., Araújo, E. M., Medeiros, V. N., Melo, T. J. A., & Pessan, L. A. of polyamide 6/bentonite clay nanocomposites. Materials Research, 15(4), 506-510. http://dx.doi.org/10.1590/S151614392012005000066. 16. Vaia, R. A., & Maguire, J. F. (2007). Polymer nanocomposites with prescribed morphology: going beyond nanoparticle-filled polymers. Chemistry of Materials, 19(11), 2736-2742. http:// dx.doi.org/10.1021/cm062693+. 17. Krishnamoorti, R., & Vaia, R. A. (2007). Polymer nanocomposites. Journal of Polymer Science. Part B, Polymer Physics, 45(24), 3252-3259. http://dx.doi.org/10.1002/polb.21319. 18. Chavarria, F., & Paul, D. R. (2006). Morphology and properties of thermoplastic polyurethane nanocomposites: Effect of organoclay structure. Polymer, 47(22), 7760-7766. http:// dx.doi.org/10.1016/j.polymer.2006.08.067. 19. Noda, N., Lee, Y.-H., Bur, A. J., Prabhu, V. M., Snyder, C. R., Roth, S. C., & McBrearty, M. (2005). Dielectric properties of nylon 6/clay nanocomposites from on-line process monitoring and off-line measurements. Polymer, 46(18), 7201-7217. http:// dx.doi.org/10.1016/j.polymer.2005.06.046. 20. Yoo, Y., & Paul, D. R. (2008). Effect of organoclay structure on morphology and properties of nanocomposites based on an amorphous polyamide. Polymer, 49(17), 3795-3804. http:// dx.doi.org/10.1016/j.polymer.2008.06.014. 21. Vlasveld, D. P. N., Groenewold, J., Bersee, H. E. N., & Picken, S. J. (2005). Moisture absorption in polyamide-6 silicate nanocomposites and its influence on the mechanical properties. Polymer, 46(26), 12567-12576. http://dx.doi.org/10.1016/j. polymer.2005.10.096. 22. Díaz, V. F. R. (2001). Preparation of organophilic clays from a Brazilian smectitic clay. Key Engineering Materials. 189191, 203-207. http://dx.doi.org/10.4028/www.scientific.net/ KEM.189-191.203. 23. Díaz, V. F. R. (1994). Preparação a nível de laboratório de algumas argilas esmectíticas organofílicas (Tese de doutorado). Departamento de Engenharia Química, Escola Politécnica, Universidade de São Paulo, São Paulo. 24. Paz, R. A., Leite, A. M. D., Araujo, E. M., Melo, T. J. A., Pessan, L. A., & Passador, F. R. (2013). Propriedades mecânicas e reológicas de nanocompósitos de poliamida 6 com argila Organofílica Nacional. Polímeros: Ciência e Tecnologia, 23(5), 682-689. http://dx.doi.org/10.4322/polimeros.2013.060. 25. Andrade, D. L. A. C. S. (2003). Desenvolvimento de nanocompósitos polipropileno/bentonita através da técnica de intercalação

60

por fusão (Dissertação de mestrado). Universidade Federal de Campina Grande, Paraíba. 26. Lewin, M., Pearce, M. E., Levon, K., Mey-Marom, A., Zammarano, M., Wilkie, C. A., & Jang, B. N. (2006). Nanocomposites at elevated temperatures: migration and structural changes. Polymers for Advanced Technologies, 17(4), 226-234. http:// dx.doi.org/10.1002/pat.684. 27. Chiu, F. C., Lai, S. M., Chen, Y. L., & Lee, T. H. (2005). Investigation on the polyamide 6/organoclay nanocomposites with or without a maleated polyolefin elastomer as a toughener. Polymer, 46(25), 11600-11609. http://dx.doi.org/10.1016/j. polymer.2005.09.077. 28. Shen, S. Z., Bateman, S., Mcmahon, P., Dell’olio, M., Gotama, J., Nguyen, T. & Yuan, Q. (2010). The effects of Clay on fire performance and thermal mechanical properties of woven glass fibre reinforced polyamide 6 nanocomposites. Composites Science and Technology, 70(14), 2063-2067. http://dx.doi. org/10.1016/j.compscitech.2010.07.027. 29. González, T. V., Salazar, C. G., Rosa, J. R., & González, V. G. (2008). Nylon 6/organoclay nanocomposites by extrusion. Journal of Applied Polymer Science, 108(5), 2923-2933. http:// dx.doi.org/10.1002/app.27307. 30. Li, T.-C., Ma, J., Wang, M., Tjiu, W. C., Liu, T., & Huang, W. (2007). Effect of Clay Addition on the Morphology and Thermal Behavior of Polyamide 6. Journal of Applied Polymer Science, 103(2), 1191-1199. http://dx.doi.org/10.1002/app.25378. 31. Oliveira, M. F. L., Oliveira, M. G., & Leite, M. C. A. M. (2011). Nanocompósitos de poliamida 6 e argila organofílica: estudo da cristalinidade e propriedades mecânicas. Polímeros: Ciência e Tecnologia, 21(1), 78-82. http://dx.doi.org/10.1590/ S0104-14282011005000015. 32. Kohan, M. I. (1973). Nylon plastics (p. 683). New York: John Wiley & Sons. 33. Yu, S., Zhao, J., Chen, G., Juay, Y. K., & Yong, M. S. (2007). The characteristics of polyamide layered-silicate nanocomposites. Journal of Materials Processing Technology, 192-193(1), 410414. http://dx.doi.org/10.1016/j.jmatprotec.2007.04.006. Received: Apr. 15, 2014 Revised: July 27, 2015 Accepted: Sept. 08, 2015

Polímeros, 26(1), 52-60, 2016


http://dx.doi.org/10.1590/0104-1428.2003

Correlation between stabilizer consumption and degree of polymerization of thermally upgraded paper aged in insulating natural ester and insulating mineral oil Larissa Mildemberger1,2, Mario Carlos Andreoli3, Guilherme Cunha da Silva1, Heloisa Nunes da Motta1, Joseane Valente Gulmine1 and Marilda Munaro1,2* Instituto de Tecnologia para o Desenvolvimento – LACTEC, Curitiba, PR, Brazil 2 Programa de Pós-graduação em Engenharia e Ciência de Materiais – PIPE, Universidade Federal do Paraná – UFPR, Curitiba, PR, Brazil 3 Companhia de Transmissão de Energia Elétrica Paulista – CTEEP, São Paulo, SP, Brazil 1

*marilda@lactec.org.br

Abstract Insulating paper holds significant importance in the insulation system of power transformers, and thus, its degradation is the subject of many studies. A successful evaluation of the degradation rate of such paper contributes to reducing downtime and avoiding equipment failure. In this work, samples of thermally upgraded paper were thermally aged in insulating natural ester (INE) and insulating mineral oil (IMO) and were evaluated by degree of polymerization (DP) and FTIR-ATR. It was possible to identify characteristic bands of dicyandiamide, an inhibitory compound of the thermal degradation of the paper, and to establish a correlation between the decrease in DP and the consumption of dicyandiamide during aging, which was observed to develop in three distinct steps for both IMO and INE. Keywords: thermally upgraded paper, thermal degradation, insulating natural ester, insulating mineral oil.

1. Introduction Transformers are essential equipment in modern electrical systems, responsible for the necessary variations in voltages for interconnection. An important part of a transformer is its insulation system, which is basically composed of an insulating fluid and an insulating solid[1]. This insulating function has been well performed by insulating mineral oil (IMO) and insulating paper for more than a century[2]. Because IMO is derived from petroleum, a non‑renewable resource, new alternatives for insulating fluids for the electrical sector have been sought. Insulating natural esters (INEs) show faster biodegradability than IMO and are derived from renewable resources; thus, INEs are strong candidates for replacing IMO[3]. Solid insulation is performed by insulating paper based on 90% cellulose coupled with hemicellulose, lignin and traces of pentose; this insulating paper is fundamentally important because it cannot be replaced while a transformer is in service, and thus, the paper determines the useful lifetime of the equipment[4]. The degree of polymerization (DP) is used to evaluate the degradation of insulation paper; it indicates the average number of α-D-glucopyranose units in a cellulose molecule. The DP of fresh paper is greater than 1000, and the end of the lifetime of a transformer is reached when the DP of the insulating paper reaches approximately 200, the moment when the paper no longer exhibits mechanical resistance and the paper’s insulating properties are compromised[5]. This DP decrease is explained by the breaking of 1.4-β glycosidic bonds in glucose monomers of cellulose chains by hydrolysis, oxidation and pyrolysis. The most significant mechanism

Polímeros, 26(1), 61-65, 2016

for cellulose degradation process in transformers is the hydrolysis. Several studies have investigated the correlation between cellulose depolymerization rate, activation energy for breaking glycosidic bonds and paper lifetime[6-9]. The production of thermally upgraded paper (TU) started in the 1950s to reduce the rate of thermal degradation of the paper[10]. Two methods for stabilizing the paper were proposed: (i) the chemical modification of the cellulose via cyanoethylation and acetylation and (ii) the addition of nitrogen compounds to protect the cellulose from oxidation – inhibitory compounds such as urea, melamine, dicyandiamide and polyacrylamide. Due to environmental concerns, chemical modification by cyanoethylation and acetylation is no longer performed[11]. By thermally accelerating the aging of kraft (KP) and thermally upgraded paper in IMO, Martins[10] observed a lower thermal degradation rate for TU than for KP, possibly due to the neutralization of the acid compounds formed during the degradation process of the insulating system by the nitrogen bases and the reduction of the effect of pyrolysis due to the reactions occurring between the inhibitory compounds and water. Authors such as McShane et al.[12], Yang et al.[13] and Frimpong et al.[14], comparing the degradation of KP in IMO and INE, observed the preservation of paper aged in INE. The authors concluded that this protection can be attributed to three main factors: (i) the INE can support hydrolysis and consume the available water inside the insulation system, reducing the potential for the hydrolysis of the paper; (ii) the INE can dry the insulating paper due to the affinity between

61

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


Mildemberger, L., Andreoli, M. C., Silva, G. C., Motta, H. N., Gulmine, J. V., & Munaro, M. the esters and water and due to the high degree of solubility of the water in this insulating fluid compared to that in IMO; and (iii) the progression of the so-called transesterification process, through which cellulose molecules are modified by the substitution of OH- groups in the cellulose by ester groups, which are larger and more stable, increasing the paper’s lifetime. However, only a few studies have been dedicated to examining the aging of thermally upgraded paper in INE. In this work, the consumption of the inhibitor compound used in upgraded paper was monitored by the FTIR-ATR technique, and a correlation between the DP of the paper aged in IMO and INE was established.

bis (ethylenediamine) for 16 h with N2 bubbling and the viscosity of this solution was measured.

2. Experimental

The calculus of dicyandiamide index was made by dividing the integral of bands between 2194-2154 cm–1, the doublet from dicyandiamide, by the integral of band of C-O bond of cellulose between 1188-914 cm–1 (Equation 1).

2.1 Materials The following materials were used: insulating thermally upgraded paper for electrical equipment from Isoeletri; insulating mineral oil AV-60IN Petrobras, with paraffin chains 48%, naphthenic chains 48% and aromatics 4%, with 17 ppm of water content; insulating natural ester soybean base Envirotemp FR3 Cargill, with 54 ppm of water content; and electrolytic copper foils.

2.2 Sample aging The paper was cut into pieces measuring 4 m in length and 15 mm in width, rolled up and dried at 100 °C for 2 hours in a vacuum oven. Electrolytic copper foils measuring 20 × 1 × 0.2 mm were prepared by first polishing them with coarse sandpaper and then with finer sandpaper. The foils were cleaned using cotton soaked in acetone and silicon carbide polishing. The copper was used in order to simulate the aging conditions of the transformer. Samples of the oil (250 mL) were placed in biochemical oxygen demand (BOD) flasks, and N2 was bubbled into each flask for 10 minutes to eliminate O2; then, a dried paper strip and the copper foil were added and the flasks were closed. The flasks were placed in an oven at 100 °C. The samples were taken at 15, 30, 33, 36, 39, 42 and 45 days of aging. After being removed from the oven, the samples were left in a dark environment for 24 hours to reach room temperature.

2.5 Dicyandiamide index FTIR analysis of the dicyandiamide content in the paper was performed by the reflectance technique, in a horizontal ATR accessory from Graseby Specac, using a rectangular ZnSe crystal measuring 10 × 50 mm. Paper samples were positioned directly on the accessory, and pressure was applied to enhance the contact between the crystal and the paper surfaces. The instrument used for analysis was a Vertex 70 Infrared Spectrophotometer from Bruker with a resolution of 4 cm–1; 32 sweeps were performed over the region between 4000 cm–1 and 650 cm–1.

int egral of the bands between dicyandiamide index =

2245 − 2040 cm −1 int egral of the band between

(1)

1188 − 914 cm −1

3. Results and discussion 3.1 Degree of polymerization Figure 1 shows the DP of the paper decreasing with aging time and there was no significant difference between the paper aged in IMO and aged in INE, indicating was not reduction in the degradation rate of the paper in INE under the conditions applied in the present work, how as suggested by some authors[12-14]. After 15 days of aging, the DP of the paper decreased but remained close to 1000, a value still considered suitable for the utilization of the paper. The DP value continued to decrease until the aging time reached between 33 and 42 days, during which the degradation rate of the paper decreased, reaching a type of plateau. During this time, the dicyandiamide most likely protected the cellulose chains, decreasing the rate of chain breakage

2.3 Oil extraction To facilitate the tests of the paper after aging, the impregnated oil was extracted via an ASE200 – Accelerated Solvent Extractor from Dionex, using hexane as solvent, at a temperature of 50 °C and pressure of 500 psi. After the extraction, the paper was dried in an oven at 75 °C for 1 hour and placed in a desiccator prior to analysis.

2.4 Degree of polymerization The DP was determined by the viscometric method according to the Brazilian standard ABNT NBR IEC 60450:2009[15], using a capillary viscometer immersed in a water bath at 20 °C±0.1 °C and a 1 mol.L–1 aqueous solution of copper(II) hydroxide bis (ethylenediamine) from Aldrich . The paper sample was milled, dried and weighted. After this, it was solved in copper(II) hydroxide 62

Figure 1. DP of the samples aged in natural ester (INE) and mineral oil (IMO) versus aging time. Polímeros, 26(1), 61-65, 2016


Correlation between stabilizer consumption and degree of polymerization of thermally upgraded paper aged in insulating natural ester and insulating mineral oil and causing the degree of polymerization to fluctuate to a lower degree. At 45 days after the beginning of the aging process, the DP was reduced to 50% of its initial value, in both IMO and INE, indicative that the mechanical properties of the material could started to be compromised. According to Shroff and Stannett[16] the reduction of 50% of the value of DP may be related to a decrease of approximately 25% of tensile strength.

3.2 Dicyandiamide index The FTIR spectrum of fresh insulating paper is shown in Figure 2, which demonstrates bands related to cellulose, the paper’s main component, as well as the characteristic bands of the inhibitor compound. The wide band at 3339 cm–1 can be attributed to the stretching vibration of the OH bond of hydroxyl groups. The band at 2918 cm–1 is due to the stretching of the C-H bond. The signal associated with the bending of CH2 bonds occurs at 1425 cm–1. The signal corresponding to the asymmetric stretching of the C-O-C bond is observed at 1159 cm–1, and that associated with the stretching of the C-O bond occurs at 1027 cm–1. The band at 896 cm–1 is possibly related to β-glycosidic bonds, which are characteristic of the cellulose molecule[17]. The bands that appear in the regions of 1658 cm–1 and 1560 cm–1 can be attributed, respectively, to the C-N bonds and to the deformation of the N-H bonds of primary amines[18]. The doublet in the 2194-2154 cm-1 region is attributed to the resonance between C≡N and C=NH, a characteristic absorption band of the dicyandiamide compound because, among the nitrogen compounds used as inhibitors in the insulating paper upgrading process (urea, guanidine, acrylonitrile, melamine and dicyandiamide), only the dicyandiamide compound shows a doublet near 2200 cm–1 in its FTIR spectrum[19]. The intensity of the dicyandiamide doublet decrease for both oils in aged paper spectrum. The inhibitor cannot be directly quantified solely by the integration of the band area of the compound’s absorption region because the ATR technique is highly affected by the contact between the sample and the crystal; therefore, the consumption of dicyandiamide was analyzed by the band ratio technique, resulting in semi-quantitative analysis (Figure 3). It can be observed that the inhibitor content in the aged paper, for both IMO and INE, decreased with aging time, indicating that the dicyandiamide was consumed during thermal aging. An increase in the consumption rate of dicyandiamide was indicated by an abrupt change in the slope of the curve beginning at 33 days of aging, corroborating the plateau in the DP behavior shown in Figure 1, with the decrease in the rate of breakage of the cellulose chains. The samples aged in INE showed slight dispersion in the data, possibly due to the fluid’s polarity and to the affinity of the dicyandiamide nitrile group for the insulating solid and liquid materials, which can promote the partition of this stabilizing agent between the paper and INE, as dicyandiamide is not chemically bound to cellulose. Again, no large differences were observed between the samples aged in IMO and INE over the same aging time, which confirmed that there was no reduction in the rate of consumption of the additive in paper aged in ester; this finding may indicate that Polímeros, 26(1), 61-65, 2016

Figure 2. FTIR-ATR spectra of thermally upgraded paper (TU), new and after 45 days aging in IMO and ENI.

Figure 3. Dicyandiamide index during aging time.

INE does not protect the paper from degradation under the effects of the inhibitor. Yang et al.[20] obtained similar results.

3.3 Correlation between DP and Dicyandiamide index Figure 4 shows the behavior of the DP in relation to the consumption of dicyandiamide. As in the study by Morais et al.[19], the inhibitor is observed to operate in three steps. The first step occurs when the paper’s DP is above 800 and the bonds that are not protected by the inhibitor can be broken, which generally begins to take place during the drying of the paper. At this stage, in the unprotected portions of the paper, acid-catalyzed hydrolysis can occur via small molecules (formic acid, acetic acid), which are readily absorbed in the paper’s structure causing oxidation of the oil and paper; this process breaks glycosidic bonds and decreases the DP. The oxidation of the paper may also occur through the formation of hydroperoxide radicals, which attack and open glucose rings, forming acids and aldehydes and also reducing the size of the cellulose chains. A certain amount of dicyandiamide is also most likely consumed through Maillard reactions[21] between primary amines (as is the case with dicyandiamide) and aldehydes (which are formed during degradation of the oil and paper). Furthermore, dicyandiamide acts as a base and reacts with 63


Mildemberger, L., Andreoli, M. C., Silva, G. C., Motta, H. N., Gulmine, J. V., & Munaro, M.

5. Acknowledgements The authors sincerely acknowledge CTEEP (Companhia de Transmissão de Energia Elétrica Paulista), CNPq (the Brazilian National Council of Scientific and Technological Development) for the benefit of the Law 8010/90, UFPR and LACTEC for their financial support.

6. References

Figure 4. Degree of polymerization (DP) as a function of dicyandiamide index.

acids in the presence of water and temperature to form ammonia and CO2, thereby consuming water. Ammonia also acts as a primary amine and reacts with aldehydes[10]. Second, the inhibitor begins to be consumed, but no significant change in DP is observed, as evidenced by the formation of a plateau between 33 and 42 days of aging. During this period, the paper is being protected while the inhibitor is being consumed by Maillard reactions and acid hydrolysis. In the third stage, approximately half of the stabilizer has already been consumed and the DP reaches one-third of its initial value at 45 days of accelerated aging. As indicated by the curves shown in Figure 4, the type of insulating fluid used did not affect the degradation processes of thermally upgraded paper in any relevant way; in fact, the effect of the stabilizing additives in the test systems was observed to be more significant. Probably when the dicyandiamide is fully consumed the paper degradation kinetics will be slower in ENI.

4. Conclusions The modification of electrical insulating paper during the upgrading process was confirmed by the FTIR-ATR technique via the identification of bands related to dicyandiamide in the regions of 1658 cm–1, 1560 cm–1 and 2194-2154 cm–1. It was possible to correlate the consumption of dicyandiamide during the paper aging process and the decrease in the DP. In addition, the paper’s behavior in natural ester was observed to be very similar to that in mineral oil. Three different regions were observed for the consumption of dicyandiamide as a function of aging, indicating that changes in the degradation mechanism took place as the inhibitor was consumed. Additionally, a slight dispersion in the data was observed for the insulating natural ester system, most likely due to the fluid’s polar nature, causing the partition of dicyandiamide between the paper and the ester. Finally, it was observed that the type of insulating fluid used exerts a minor effect on the degradation of thermally upgraded paper while the dicyandiamide is not fully consumed. 64

1. Fofana, I., Bouslimi, H., Hemmatjou, C., Volat, K., & Tahiri, K. (2014). Relationship between static electrification of transformer oils with turbidity and spectrophotometry measurements. International Journal of Electrical Power & Energy Systems, 54, 38-44. http://dx.doi.org/10.1016/j.ijepes.2013.06.037. 2. Prevost, T. A. (2005). Thermally upgraded insulation in transformer. In Proceedings of Electrical Insulation Conference and Electrical Manufacturing Expo (pp. 120-125). Indianapolis: IEEE. 3. Ciuriuc, A., Vihacencu, M. S., Dumitran, L. M., & Notingher, P. V. (2012). Comparative study on power transformers vegetable oil and mineral oil ageing. In Proceedings of International Conference on Applied and Theoretical Electricity (pp. 1-6). Craiova: IEEE. 4. Martins, M. A. G. (2007). Furfuraldeído: um indicador prático da degradação térmica do papel kraft de transformadores. Ciência e Tecnologia dos Materiais, 19(1-2), 25-33. 5. Liland, K. B., Ese, M. H. G., Selsbak, C. M., & Lundgaard, L. (2011). Ageing of oil impregnated thermally upgraded papers. In Proceedings of IEEE International Conference on Dielectric Liquids (pp. 1-5). Trondheim: IEEE. 6. Jalbert, J., Gilbert, R., Tétreault, P., Morin, B., & LessardDéziel, D. (2007). Identification of a chemical indicator of the rupture of 1,4-beta-glycosidic bonds of cellulose in an oil-impregnated insulating paper system. Cellulose, 14(4), 295-309. http://dx.doi.org/10.1007/s10570-007-9124-1. 7. Gilbert, R., Jalbert, J., Tétreault, P., Morin, B., & Denos, Y. (2009). Kinetics of the production of chain-end groups and methanol from the depolymerization of cellulose during the ageing of paper/oil systems. Part 1: standard wood kraft insulation. Cellulose, 16(2), 327-338. http://dx.doi.org/10.1007/ s10570-008-9261-1. 8. Gilbert, R., Jalbert, J., Duchesne, S., Tetreault, P., Morin, B., & Denos, Y. (2010). Kinetics of the production of chain-end groups and methanol from the depolymerization of cellulose during the ageing of paper/oil systems. Part 2: thermallyupgraded insulating papers. Cellulose, 17(2), 253-269. http:// dx.doi.org/10.1007/s10570-009-9365-2. 9. Setnescu, R., Badicu, L. V., Dumitran, L. M., Notingher, P. V., & Setnescu, T. (2014). Thermal lifetime of cellulose insulation material evaluated by an activation energy based method. Cellulose, 21(1), 823-833. http://dx.doi.org/10.1007/ s10570-013-0087-0. 10. Martins, M. A. (2007). Monitorização da degradação térmica do papel isolante usado em transformadores: papel “thermally upgraded” versus papel kraft. Ciência e Tecnologia dos Materiais, 19(1-2), 14-18. 11. Prevost, T. A. (2009). Dielectric properties of natural esters and their influence on transformer insulation system design and performance: an update. In Proceedings of Power & Energy Society General Meeting (pp. 1-7), Calgary: IEEE. 12. McShane, C. P., Rapp, K. J., Corkran, J. L., Gauger, G. A., & Luksich, J. (2001). Aging of paper insulation in natural ester dielectric fluid. In Proceedings of Transmission and Distribution Conference and Exposition (Vol. 2, pp. 675-679). Atlanta: IEEE. Polímeros, 26(1), 61-65, 2016


Correlation between stabilizer consumption and degree of polymerization of thermally upgraded paper aged in insulating natural ester and insulating mineral oil 13. Yang, L., Liao, R., Caixin, S., & Zhu, M. (2011). Influence of vegetable oil on the thermal aging of transformer paper and its mechanism. IEEE Transactions on Dielectrics and Electrical Insulation, 18(3), 692-700. http://dx.doi.org/10.1109/ TDEI.2011.5931054. 14. Frimpong, G. K., Oommen, T. V., & Asano, R. (2011). A survey of aging characteristics of cellulose insulation in natural ester and mineral oil. IEEE Electrical Insulation Magazine, 27(5), 36-48. http://dx.doi.org/10.1109/MEI.2011.6025367. 15. Associação Brasileira de Normas Técnicas – ABNT. (2009). ABNT NBR IEC 60450: medição do grau de polimerização viscosimétrico médio de materiais celulósicos novos e envelhecidos para isolação elétrica. Rio de Janeiro. 16. Shroff, D. H., & Stannett, A. W. (1985). A review of paper aging in power transformers. IEE Proceedings. Part C. Generation, Transmission and Distribution, 132(6), 312-319. http://dx.doi. org/10.1049/ip-c.1985.0052. 17. Yu, X., Tong, S., Ge, M., Wu, L., Zuo, J., Cao, C., & Song, W. (2013). Synthesis and characterization of multi-aminofunctionalized cellulose for arsenic adsorption. Carbohydrate Polymers, 92(1), 380-387. http://dx.doi.org/10.1016/j. carbpol.2012.09.050. PMid:23218309.

Polímeros, 26(1), 61-65, 2016

18. Urreaga, J. M., & Orden, M. U. (2007). Modification of cellulose with amino compounds: a fluorescence study. Carbohydrate Polymers, 69(1), 14-19. http://dx.doi.org/10.1016/j. carbpol.2006.08.019. 19. Morais, R. M., Mannheimer, W. A., Carballeira, M., & Noualhaguet, J. C. (1999). Furfural analysis for assessing degradation of thermally upgraded papers in transformer insulation. IEEE Transactions on Dielectrics and Electrical Insulation, 6(2), 159-163. http://dx.doi.org/10.1109/94.765905. 20. Yang, L., Liao, R., Sun, C., & Sun, H. (2008). Study on the influence of natural ester on thermal ageing characteristics of oil-paper in power transformer. In Proceedings of International Conference on High Voltage Engineering and Application (pp. 437-440). Chongqing: IEEE. 21. Lundgaard, L. E., Hansen, W., Linhjell, D., & Painter, T. J. (2004). Aging of oil-impregnated paper in power transformers. IEEE Transactions on Power Delivery, 19(1), 230-239. http:// dx.doi.org/10.1109/TPWRD.2003.820175. Received: Nov. 14, 2014 Revised: June 25, 2015 Accepted: Sept. 11, 2015

65


http://dx.doi.org/10.1590/0104-1428.2123

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

Poly(lactic acid)/thermoplastic starch sheets: effect of adipate esters on the morphological, mechanical and barrier properties Marianne Ayumi Shirai1,2*, Juliana Bonametti Olivato1, Ivo Mottin Demiate3, Carmen Maria Olivera Müller4, Maria Victória Eiras Grossmann1 and Fabio Yamashita1 Department of Food Science and Technology, Universidade Estadual de Londrina – UEL, Londrina, PR, Brazil 2 Department of Food Technology, Universidade Tecnológica Federal do Paraná – UTFPR, Campus Londrina, Londrina, PR, Brazil 3 Department of Food Engineering, Universidade Estadual de Ponta Grossa – UEPG, Ponta Grossa, PR, Brazil 4 Department of Food Science and Technology, Universidade Federal de Santa Catarina – UFSC, Florianópolis, SC, Brazil 1

*marianneshirai@utfpr.edu.br

Abstract Blends of poly(lactic acid) (PLA) and thermoplastic starch (TPS) plasticized with adipate esters (diisodecyl adipate and diethyl adipate) having different molecular weight were used to produce sheets. The calendering-extrusion process at a pilot scale was used, and the mechanical, barrier, and morphological characterization of the obtained materials were performed. The increase in the TPS content affected the mechanical properties of the sheets by increasing the elongation and decreasing the rigidity. TPS conferred a more hydrophilic character to the sheets, as observed from the water vapor permeability results. The sheets plasticized with diisodecyl adipate (DIA), having a higher molecular weight, had better mechanical and barrier properties than diethyl adipate (DEA) plasticized sheets, indicating that DIA was more effective as plasticizer. Micrographs obtained by confocal laser microscopy and scanning electron microscopy showed different morphologies when different proportions of PLA and TPS were used (dispersed or co-continuous structures), which were strongly associated with the mechanical and barrier properties. Keywords: biodegradable material, polymeric blend, plasticizer, mechanical properties.

1. Introduction The interest in producing plastic materials from natural resources is considerably increasing as the need for the reduction of the amount of plastic waste in the environment becomes urgent[1]. Due to its biodegradability, low cost, and worldwide availability, starch in the form of thermoplastic starch (TPS) has been extensively studied as a main component in the production of biodegradable materials. TPS is usually obtained by destroying the crystalline structure of native starch through an extrusion process in the presence of plasticizers, such as glycerol[2-4]. Unfortunately, TPS-based materials are hygroscopic and have limited performance. To overcome this deficiency, the TPS must be blended with another biodegradable polymer to produce materials for packaging and industrial applications[5,6]. In this context, the blend of TPS with poly (lactic acid) (PLA) is promising because, in addition to being compostable, PLA is produced from a renewable resource. PLA has received much attention as the most innovative alternative to conventional petroleum-based polymers and has been intensively studied due to its environmentally friendly characteristics such as biocompatibility, sustainability, and potentially useful physical and mechanical properties.

66

The addition of TPS into a PLA matrix can also decrease the material cost and increase its biodegradation rate[5,7,8]. Blending native starch with PLA could increase its rigidity but simultaneously greatly reduce its elongation at break and impact strength[9]. Many factors are believed to cause the deterioration in mechanical properties of PLA/starch blends. First, PLA has an inherent brittleness as evidenced by its relatively low tensile strain at break, and toughness[10,11]. Another explanation is linked to the inherent incompatibility between the hydrophilic starch and the hydrophobic PLA that affects the interfacial adhesion[5,7,12-15]. It is well known that the mechanical and barrier properties of polymer blends are strongly related to their morphology. The relationship of the mechanical and water vapor barrier properties to the morphology of PLA and starch blends has been studied by other authors[16-18]. Blends of TPS and PLA can result in a wide variety of morphologies from a dispersed starch/PLA matrix, a co-continuous starch/PLA structure, to inverted phase morphologies, in which the starch is the continuous phase[19]. Addition of an adequate plasticizer and compatibilizer can improve the processability and flexibility of PLA, enabling the blending with starch to produce sheets with adequate

Polímeros, 26(1), 66-73, 2016


Poly(lactic acid)/thermoplastic starch sheets: effect of adipate esters on the morphological, mechanical and barrier properties mechanical properties. Another way to enhance the interfacial affinity is the addition of previously gelatinized starch to the blend. The disintegrated starch granules overcome the strong interactions between the starch molecules and, in the presence of water and other plasticizers, leads to good dispersion[20,21]. In the previous work adipate esters, compared with citrate esters, improved the mechanical properties of TPS/PLA films produced by blown extrusion[22] and TPS/PLA sheets obtained by calendering extrusion,[23] reducing the elastic modulus and the tensile strength and increasing the elongation at rupture. Adipate esters were also studied by other authors[24,25] to enable the production of flexible PLA films. The current study focuses on the production of sheets with different proportions of TPS and PLA and plasticized with two different adipate esters. The morphologies of these sheets are correlated with the mechanical and barrier properties.

2. Materials and Methods 2.1 Materials For the production of biodegradable sheets, PLA Ingeo 4043D (Natureworks LLC, Cargill, Blair, Nebraska, USA) with average molecular weight of 200 kDa and native cassava starch (Indemil, Paranavaí, Brazil) were used. Commercial glycerol (Dinamica, Diadema, Brazil) was used as a plasticizer for the starch. Diisodecyl adipate (426.67 g/mol) and diethyl adipate (202.25 g/mol) (Sigma Aldrich, Steinheim, Germany) were used as plasticizing agents for the PLA.

2.2 PLA/TPS Sheet Production Biodegradable sheets with different proportions of PLA and TPS (60/40, 50/50, 40/60, 30/70, 20/80 and 10/90, w/w) were produced, and the concentration of glycerol was 33 g/100 g starch. Two different adipate esters, diisodecyl adipate (DIA) and diethyl adipate (DEA) were used, and the concentration employed was 10 g adipate ester/100 g PLA. All of the components (PLA, starch, glycerol, and adipate ester) were manually mixed and processed using a pilot twin‑screw extruder (BGM, D-20 model, Brazil) with a screw diameter of 20 mm (L/D = 35). The screw speed was 100 rpm, the feed speed was 30 rpm, and a temperature profile from entrance to exit of 100/180/180/180/180°C was used. The extruded cylindrical profiles obtained were pelletized and extruded again in the same twin-screw extruder coupled with a calender (AX-Plasticos, Brazil) for sheet production. The temperature profile from entrance to exit was 100/170/170/170/175°C, the distance between the calender rolls was 0.8 mm, and the roll speed was adjusted for each formulation to maintain a continuous process.

2.3 Thickness The film thicknesses were measured using a digital micrometer (Starrett, Brazil) at ten different locations on the films. Polímeros, 26(1), 66-73, 2016

2.4 Mechanical properties Tensile tests were performed using a texture analyzer (Stable Micro Systems, TA XTplus, England) based on the ASTM D882-02 standard[26]. The specimens were previously conditioned at 23 ± 2°C and 53 ± 2% for 48 hours, and the tensile strength (MPa), elongation at break (%), and Young’s modulus (MPa) were measured. The tests were repeated ten times for each formulation.

2.5 Water vapor permeability The water vapor permeability (WVP) of the sheets was determined gravimetrically, according to the ASTM E96-00 standard[27]. The sheets were previously stored at 23 ± 2°C and 53 ± 2% for 48 hours and then fixed in the 60-mm-diameter cylindrical aluminum cells. The interior of the cell was filled with calcium chloride (0% RH), and the device was stored at 25°C in a desiccator containing a saturated sodium chloride solution (75% RH). The samples were weighed in semi-analytical balance (Marte, Brazil) every 12 h during the 5 days of testing. The mass gain (m) of the cell was plotted as a function of the time (t). The slope of the line was calculated using linear regression (R2>0.99), and the water vapor permeation ratio (WVPR) was calculated using Equation 1: WVPR = ( m / t ) x (1/ A )

(1)

where m/t is the angular coefficient of the curve and A is the sample permeation area (m2). The WVP (g/m.day.Pa) was calculated as follows (Equation 2): WVP = (WVPR x e ) / p ( RH 1 – RH 2 )

(2)

where e is the mean sample thickness (m), p is the water vapor saturation pressure at the assay temperature (Pa), RH1 is the relative humidity of the desiccator and RH2 is the relative humidity in the interior of the permeation cell. The test was conducted in duplicate.

2.6 Morphological characterization The sheet fracture morphology was evaluated using scanning electron microscopy (Philips, FEI Quanta 200 model, Japan). The sheets were fractured after immersion in liquid nitrogen and gold coated using a sputter coater (Bal-Tec, SCD-050 model, Balzers, Liechtenstein). All of the samples were examined using an accelerating voltage of 20 kV at a magnification of 400x. A confocal laser scanning microscope (Olympus LEXT, OLS400 model, Germany) was used to obtain surface images of the sheets.

2.7 Statistical analysis The obtained results were evaluated using the analysis of variance (ANOVA), and the treatment means were compared using Tukey’s test at a 5% significance level (p<0.05) with Statistica 7.0 software (STATSOFT, USA). 67


Shirai, M. A., Olivato, J. B., Demiate, I. M., Müller, C. M. O., Grossmann, M. V. E., & Yamashita, F.

3. Results and Discussion 3.1 Thickness It was not possible to produce sheets with PLA/TPS proportions of 20/80 and 10/90, because the extruded materials were too stiff to be processed through the calender rolls. The other formulations had good processability, and the sheets showed slightly rough surfaces. As the starch amount increased, the color became more yellowish. The thickness of the PLA/TPS sheets with adipate esters added ranged from 431 to 815 µm (Table 1). Considering the effect of TPS addition, the increase in its concentration conferred thicker sheets for PLA plasticized with DIA. In  the same way, evaluating the effect of the plasticizer type, the sheets with DIA were less thick than those with DEA for the same TPS/PLA compositions. This result could be related to the DIA plasticizing effect that enhanced the flexibility and allowed increased stretching of the sheets by the calender rolls. During the calendering extrusion process, the sheet thickness was controlled by the distance between the rolls, the speed of the rolls, and the stretching capacity of the formulation, as observed in our previous study[23].

3.2 Morphological characterization The fracture micrographs of the PLA/TPS sheets examined using SEM at 400x magnification are presented in Figure 1. For the 60/40 and 50/50 PLA/TPS sheets with DEA and DIA, starch granules with different sizes were dispersed in the PLA matrix. When the concentration was reduced to 30/70 PLA/TPS, PLA was present as a dispersed phase in the TPS matrix, evidencing a phase inversion point. According to the literature,[28] the co-continuity occurs at the phase inversion point and when both polymers are approximately in equal concentration. Furthermore, in the case of starch/PLA blends, co-continuity occurs for a low proportion of starch because the starch phase is less viscous and has a high interfacial tension. Generally, for up to 40/60 PLA/TPS, a dispersed starch phase in the matrix morphology was observed. Below this proportion of PLA, a co-continuous morphology was obtained. A similar behavior was also observed in blends of PLA and starch with varied amylose contents[16] and in blends of PLA and chemically modified corn starch[18]. In the sheets with higher proportions of PLA, some spaces between the PLA and TPS phases were evident, showing the interface between the polymers. This result could be associated with poor interaction and incompatibility Table 1. Thickness of PLA/TPS sheets plasticized with adipate ester. PLA/TPS 60/40 50/50 40/60 30/70

Thickness (µm) DEA DIA 764 ± 2b 538 ± 7d 749 ± 8b 431 ± 69e b 753 ± 9 619 ± 5c 815 ± 21a 736 ± 31b

Different letters represent significant differences (p ≤ 0.05) between the means obtained by Tukey’s test. PLA = poly(lactic acid); TPS = thermoplastic starch; DEA = diethyl adipate; DIA = diisodecyl adipate.

68

between the starch and PLA because of the difference in hydrophobicity, as observed in other reports[7,12-14,17,18,29-31]. Confocal laser microscopy (CLM) is usually employed to study the surface morphology of polymeric films and sheets and confer information about the surface topography and heterogeneity. In this study, CLM was used to analyze the surface properties of the PLA/TPS sheets plasticized with DEA or DIA, and the obtained images are shown in Figure 2. Sheets with a higher proportion of PLA had increase surface roughness. This result can be attributed to the starch granules dispersed in the PLA matrix, as observed in the fractured SEM images. The roughness decreased as the PLA concentration decreased due phase inversion, in which TPS was the continuous phase. CLM was also used to evaluate the surface roughness of the starch films treated by plasma[1] and to identify the blend compositions that correspond to the beginning of partial continuity until total continuity of the poly(butylene succinate-co-adipate) (PBSA) and TPS phases[32].

3.3 Mechanical properties The Figure 3 shows the mechanical properties in terms of the tensile strength (MPa), elongation at rupture (%) and Young´s modulus (MPa) of the PLA/TPS sheets plasticized with adipate esters. As the TPS concentration increased, the tensile strength and Young´s modulus of all of the sheets significantly decreases. For the elongation at break, only the sheets plasticized with DEA exhibited a significant increase when the TPS concentration increased. The sheets plasticized with DIA maintained this property for all of the tested TPS concentrations. For the blends with PLA as a continuous phase, DIA interacted better with it and aided in the induced crystallization due chain alignment during the tensile test. This result caused tensile energy diffusion and led to plastic deformation of the PLA matrix. The statistically equal mechanical properties for the 30/70 PLA/TPS sheets support these results. Different results were observed in our previous study[23] in which DEA with lower molecular weight promoted greater values of elongation of break than DIA. Xiong et al.[21] and Xiong et al.[33] evaluated the effects of plant oils such as Tung oil anhydride and epoxidized soybean oil as plasticizers in PLA/TPS blends containing approximately 60% of PLA and reported a decrease in the tensile strength and an increase in the elongation at break of the blends. Comparing the effects of the adipate esters as plasticizers, DIA promoted higher values of the elongation at break and the Young´s modulus; however, the same tensile strength values were observed for the two plasticizers. It is possible to conclude that adipate esters with higher molecular weight promoted materials with greater elongation, as observed by other authors[24] that used polyadipate as a plasticizer for PLA films. As observed in the SEM and CLM images, the starch granules dispersed in the PLA matrix act as defects and lead to serious stress concentration, which the material easier to fracture[21]. The presence of spaces between the starch and PLA also impaired the load transfer under stress. Therefore, the tensile strength and elongation at break of Polímeros, 26(1), 66-73, 2016


Poly(lactic acid)/thermoplastic starch sheets: effect of adipate esters on the morphological, mechanical and barrier properties

Figure 1. SEM (400 X) of PLA/TPS sheets plasticized with adipate esters. PolĂ­meros, 26(1), 66-73, 2016

69


Shirai, M. A., Olivato, J. B., Demiate, I. M., Müller, C. M. O., Grossmann, M. V. E., & Yamashita, F.

Figure 2. Confocal microscopy images of PLA/TPS sheets plasticized with adipate esters.

the materials are directly linked to interfacial adhesion between the polymers used in the blends. Similar results for PLA/starch blends plasticized with acetyl triethyl citrate were reported[12].

produced by Teixeira et al.[34] due the addition of cassava bagasse.

The mechanical properties of the 30/70 PLA/TPS samples were similar for both plasticizers, suggesting that in the case of a dispersed phase (PLA) surrounded by a continuous phase (TPS), the mechanical properties of these immiscible blends depend on those of the continuous phase because it absorbs the stress and energy when the material is under load[28]. When percolation occurs, a part or total of one phase forms a continuous structure that permeates the entire sample and eventually dominates the properties of the blend. In our case, as the TPS levels increased, the interface properties improved, as observed by MEV. As a result, the tensile strength of the sheets was expected to increase; however, this result did not occur because the TPS has a lower tensile strength and Young’s modulus than the PLA.

The WVP of the PLA/TPS sheets with added adipate esters ranged from 2.94 x 10-6 to 6.98 x 10-6 g day-1 m-1 Pa-1, as shown in Figure 4. The values obtained are lower than those observed for cassava starch films[35,36] and PLA/TPS films[29] and greater than PLA films with synthetic phenolic antioxidants added[37]. The presence of PLA promoted a hydrophobic character of the material and was responsible for decreasing the WVP because the material permeability is strongly influenced by the hydrophobic or hydrophilic nature of its components.

Comparing the obtained results with other results reported in the literature, lower values of the tensile strength, Young’s modulus and elongation at break were obtained in PLA/TPS (30 g PLA / 100 g TPS) films produced by extrusion and thermopressing[29,30] In this case, because TPS was the continuous phase, it was responsible for the mechanical properties of the films, resulting in more fragile materials. PLA/TPS films (20% TPS) with increased resistance were 70

3.4 Water vapor permeability (WVP)

The increase in the TPS concentration (from 60/40 to 30/70) elevated the WVP because of the hydrophilic character of the starch. Furthermore, at high starch contents, water could easily saturate the surface of the blend, penetrate into these voids and be quickly absorbed by the starch, resulting in higher water permeation[16]. Comparing the type of adipate ester, the sheets with DIA had lower WVP than sheets with DEA, mostly in the materials with a higher proportion of PLA. To support the discussion of the results, the plasticizing capacity of the adipate esters was calculated by dividing the molecular weight by the hydrogen bond capacity of each component, resulting Polímeros, 26(1), 66-73, 2016


Poly(lactic acid)/thermoplastic starch sheets: effect of adipate esters on the morphological, mechanical and barrier properties

Figure 4. Water vapor permeability of PLA/TPS sheets plasticized with DEA and DIA. Different letters represent significant difference (p ≤ 0.05) between the means obtained by Tukey’s test.

4. Conclusion Calendering extrusion is a feasible way to produce PLA/TPS sheets at a pilot scale with great potential to be used in packaging production. Different proportions of PLA and TPS in the blends confer sheets with modified morphological structures (dispersed or co-continuous structures), which significantly affect the mechanical and barrier properties. Adipate esters are excellent plasticizers for PLA in PLA/TPS sheets. DIA, because of its higher molecular weight, was the better plasticizer for PLA conferring sheets with higher elongation at break, tensile strength, Young´s modulus, as well as reduced WPV.

5. Acknowledgements The authors would like to thank CAPES, CNPq and Fundação Araucaria for their financial support and research grants and the Microscopy Laboratory of UEL for the use of their scanning electron microscope.

6. References Figure 3. Mechanical properties of PLA/TPS sheets plasticized with DEA and DIA. Different letters represent significant difference (p ≤ 0.05) between the means obtained by Tukey’s test.

in 9.9 x 103 and 4.7 x 103 for DEA and DIA, respectively. DEA presented a higher value and consequently a better plasticizing capacity that increased the intermolecular distance between the chains and the free volume, facilitating the penetration of water molecules into the polymer matrix. This fact and the longer chains of the DIA explain the lower WVP values of the DIA-added PLA/TPS sheets. Similar results for PLA/TPS sheets plasticized with DEA were observed in our previous study[23]. Finally, the final WVP is a balance between the hydrophilic and hydrophobic character of the blend components and the interface properties. The greater the tension at the interface (less compatible blends), the greater the spaces between the phases, which facilitates water diffusion. Polímeros, 26(1), 66-73, 2016

1. Bastos, D. C., Santos, A. E. F., Silva, M. L. V. J., & Simão, R. A. (2009). Hydrophobic corn starch thermoplastic films produced by plasma treatment. Ultramicroscopy, 109(8), 1089-1093. http://dx.doi.org/10.1016/j.ultramic.2009.03.031. PMid:19345017. 2. Avérous, L., Fringant, C., & Moro, L. (2001). Plasticized starch-cellulose interaction in polysaccharide composites. Polymer, 42(15), 6565-6572. http://dx.doi.org/10.1016/S00323861(01)00125-2. 3. Liu, H., Xie, F., Yu, L., Chen, L., & Li, L. (2009). Thermal processing of starch-based polymers. Progress in Polymer Science, 34(12), 1348-1368. http://dx.doi.org/10.1016/j. progpolymsci.2009.07.001. 4. Zullo, R., & Iannace, S. (2009). The effects of different starch sources and plasticizers on film blowing of thermoplastic starch: Correlation among process, elongational properties and macromolecular structure. Carbohydrate Polymers, 77(2), 376-383. http://dx.doi.org/10.1016/j.carbpol.2009.01.007. 5. Li, H., & Huneault, M. A. (2011). Comparison of sorbitol and glycerol as plasticizers for thermoplastic starch in TPS/PLA blends. Journal of Applied Polymer Science, 119(4), 2439-2448. http://dx.doi.org/10.1002/app.32956. 71


Shirai, M. A., Olivato, J. B., Demiate, I. M., Müller, C. M. O., Grossmann, M. V. E., & Yamashita, F. 6. Abdillahi, H., Chabrat, E., Rouilly, A., & Rigal, L. (2013). Influence of citric acid on thermoplastic wheat flour / poly(lactic acid) blends. II. Barrier properties and water vapour sorption isotherms. Industrial Crops and Products, 50, 104-111. http:// dx.doi.org/10.1016/j.indcrop.2013.06.028. 7. Martin, O., & Averous, L. (2001). Poly(lactic acid): plasticization and properties of biodegradable multiphase systems. Polymer, 42(14), 6209-6219. http://dx.doi.org/10.1016/S00323861(01)00086-6. 8. Babu, R. P., O’Connor, K., & Seeram, R. (2013). Current progress on bio-based polymers and their future trends. Progress in Biomaterials, 2(8), 1-16. http://dx.doi.org/10.1186/21940517-2-8. 9. Jacobsen, S., & Fritz, H. G. (1999). Plasticizing polylactide: The effect of different plasticizers on the mechanical properties. Polymer Engineering and Science, 39(7), 1303-1310. http:// dx.doi.org/10.1002/pen.11517. 10. Bhardwaj, R., & Mohanty, A. K. (2007). Advances in the properties of polylactides based materials: a review. Journal of Biobased Materials Bioenergy, 1(2), 191-209. http://dx.doi. org/10.1166/jbmb.2007.023. 11. Anderson, K. S., Schreck, K. M., & Hillmyer, M. (2008). Toughening polylactide. Polymer Reviews, 48(1), 85-108. http://dx.doi.org/10.1080/15583720701834216. 12. Zhang, J. F., & Sun, X. Z. (2004). Mechanical properties of poly(lactic acid)/starch composites compatibilized by maleic anhydride. Biomacromolecules, 5(4), 1446-1451. http://dx.doi. org/10.1021/bm0400022. PMid:15244463. 13. Wang, N., Yu, J., Chang, P. R., & Ma, X. (2008). Influence of formamide and water on the properties of thermoplastic starch/poly(lactic acid) blends. Carbohydrate Polymers, 71(1), 109-118. http://dx.doi.org/10.1016/j.carbpol.2007.05.025. 14. Kozlowski, M., Masirek, R., Piorkowska, M., & GazickiLipman, M. (2007). Biodegradable blends of poly(L-lactide) and starch. Journal of Applied Polymer Science, 105(1), 269277. http://dx.doi.org/10.1002/app.26088. 15. Lim, L. T., Auras, R., & Rubino, M. (2008). Processing technologies for poly(lactic acid). Progress in Polymer Science, 33(8), 820852. http://dx.doi.org/10.1016/j.progpolymsci.2008.05.004. 16. Ke, T., Sun, S. X., & Seib, P. (2003). Blending of Poly(lactic acid) and starches containing varying amylose content. Journal of Applied Polymer Science, 89(13), 3639-3646. http://dx.doi. org/10.1002/app.12617. 17. Jang, W. Y., Shin, B. Y., Lee, T. J., & Narayan, R. (2007). Thermal properties and morphology of biodegradable PLA/starch compatibilized blends. Journal of Industrial and Engineering Chemistry, 13(3), 457-464. Retrieved in 22 February 2015, from http://infosys.korea.ac.kr/research/tech/periodicals/view. php?seq=581012 18. Shin, B. Y., Jang, S. H., & Kim, B. S. (2011). Thermal, morphological, and mechanical properties of biobased and biodegradable blends of poly(lactic acid) and chemically modified thermoplastic starch. Polymer Engineering and Science, 51(5), 826-834. http://dx.doi.org/10.1002/pen.21896. 19. Huneault, M. A., & Li, H. (2007). Morphology and properties of compatibilized polylactide/thermoplastic starch blends. Polymer, 48(1), 270-280. http://dx.doi.org/10.1016/j.polymer.2006.11.023. 20. Park, J. W., Im, S. S., Kim, S. H., & Kim, Y. H. (2000). Biodegradable polymer blends of poly(L-lactic acid) and gelatinized starch. Polymer Engineering and Science, 40(12), 2539-2550. http://dx.doi.org/10.1002/pen.11384. 21. Xiong, Z., Yang, Y., Feng, J., Zhang, X., Zhang, C., Tang, Z., & Zhu, J. (2013). Preparation and characterization of poly(lactic acid)/ starch composites toughened with epoxidized soybean oil. Carbohydrate Polymers, 92(1), 810-816. http://dx.doi. org/10.1016/j.carbpol.2012.09.007. PMid:23218370. 72

22. Shirai, M. A., Grossmann, M. V. E., Mali, S., Yamashita, F., Garcia, P. S., & Müller, C. M. O. (2013). Development of biodegradable flexible films of starch and poly(lactic acid) plasticizes with adipate or citrate esters. Carbohydrate Polymers, 92(1), 19-22. http://dx.doi.org/10.1016/j.carbpol.2012.09.038. PMid:23218260. 23. Shirai, M. A., Müller, C. M. O., Grossmann, M. V. E., & Yamashita, F. (2015). Adipate and citrate esters as plasticizers for poly(lactic acid) / thermoplastic starch sheets. Journal of Polymer and Environment, 23(1), 54-61. http://dx.doi. org/10.1007/s10924-014-0680-9. 24. Martino, V. P., Ruseckaite, R. A., Jiménez, A., & Averous, L. (2010). Correlation between composition, structure and properties of poly(lactic acid)/polyadipate-based nano-biocomposites. Macromolecular Materials and Engineering, 295(6), 551-558. http://dx.doi.org/10.1002/mame.200900351. 25. Martino, V. P., Jiménez, A., & Ruseckaite, R. A. (2009). Processing and characterization of poly(lactic acid) films plasticized with commercial adipates. Journal of Applied Polymer Science, 112(4), 2010-2018. http://dx.doi.org/10.1002/ app.29784. 26. American Society for Testing and Material – ASTM. (2002). D882-02: Standard test methods for tensile properties of thin plastic sheeting. Philadelphia: ASTM. 27. American Society for Testing and Material – ASTM. (1996). E96-00: Standard test methods for water vapor transmission of materials. Philadelphia: ASTM. 28. Schwach, E., & Avérous, L. (2004). Starch-based biodegradable blends: morphology and interface properties. Polymer International, 53(12), 2115-2124. http://dx.doi.org/10.1002/ pi.1636. 29. Müller, C. M. O., Pires, A. T. N., & Yamashita, F. (2012). Characterization of thermoplastic starch/poly(lactic acid) blends obtained by extrusion and thermopressing. Journal of the Brazilian Chemical Society, 23(3), 426-434. http://dx.doi. org/10.1590/S0103-50532012000300008. 30. Soares, F. C., Yamashita, F., Müller, C. M. O., & Pires, A. T. N. (2014). Effect of cooling and coating on thermoplastic starch/ poly(lactic acid) blend sheets. Polymer Testing, 33, 34-39. http://dx.doi.org/10.1016/j.polymertesting.2013.11.001. 31. Karagoz, S., & Ozkoc, G. (2013). Effects of diisocyanate compatibilizer on the properties of citric acid modified thermoplastic starch / poly(lactic acid) blends. Polymer Engineering and Science, 53(10), 2183-2193. http://dx.doi. org/10.1002/pen.23478. 32. Khalil, F., Galland, S., Cottaz, A., Joly, C., & Degraeve, P. (2014). Polybutylene succinate adipate/starch blends: A morphological study for the design of controlled release films. Carbohydrate Polymers, 108, 271-280. http://dx.doi. org/10.1016/j.carbpol.2014.02.062. PMid:24751274. 33. Xiong, Z., Li, C., Ma, S., Feng, J., Yang, Y., Zhang, R., & Zhu, J. (2013). The properties of poly(lactic acid)/starch blends with a functionalized plant oil: Tung oil anhydride. Carbohydrate Polymers, 95(1), 77-84. http://dx.doi.org/10.1016/j. carbpol.2013.02.054. PMid:23618242. 34. Teixeira, E. M., Curvelo, A. A. S., Corrêa, A. C., Marconcini, J. M., Glenn, G. M., & Mattoso, L. H. C. (2012). Properties of thermoplastic starch from cassava bagasse and cassava starch and their blends with poly(lactic acid). Industrial Crops and Products, 37(1), 61-68. http://dx.doi.org/10.1016/j. indcrop.2011.11.036. 35. Alves, V. D., Mali, S., Beléia, A., & Grossmann, M. V. E. (2007). Effect of glycerol and amylose enrichment on cassava starch film properties. Journal of Food Engineering, 78(3), 941-946. http://dx.doi.org/10.1016/j.jfoodeng.2005.12.007. 36. Mali, S., Sakanaka, L. S., Yamashita, F., & Grossmann, M. V. E. (2005). Water sorption and mechanical properties of Polímeros, 26(1), 66-73, 2016


Poly(lactic acid)/thermoplastic starch sheets: effect of adipate esters on the morphological, mechanical and barrier properties cassava starch films and their relation to plasticizing effect. Carbohydrate Polymers, 60(3), 283-289. http://dx.doi. org/10.1016/j.carbpol.2005.01.003. 37. Jamshidian, M., Tehrany, E. A., Cleymand, F., Leconte, S., Falher, T., & Desobry, S. (2012). Effects of synthetic phenolic antioxidants on physical, structural, mechanical and barrier

PolĂ­meros, 26(1), 66-73, 2016

properties of poly lactic acid film. Carbohydrate Polymers, 87(2), 1763-1773. http://dx.doi.org/10.1016/j.carbpol.2011.09.089. Received: Feb. 22, 2015 Revised: July 10, 2015 Accepted: Sept. 14, 2015

73


http://dx.doi.org/10.1590/0104-1428.2033

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

TLC/IR (UATR) off-line coupling for the characterization of additives in EPDM rubber compositions Denis Damazio1,2, Eunice Aparecida Campos1,3, Milton Faria Diniz3, Elizabeth da Costa Mattos1,3 and Rita de Cássia Lazzarini Dutra1* Instituto Tecnológico de Aeronáutica – ITA, São José dos Campos, SP, Brazil 2 Zanaflex Compostos de Borracha, Cotia, SP, Brazil 3 Divisão de Química – AQI, Instituto de Aeronáutica e Espaço – IAE, São José dos Campos, SP, Brazil 1

*ritacld@ita.br

Abstract The knowledge of the components that constitutes a rubber composition is important to justify the properties of the final device, particularly when it comes to elastomeric compositions used in the aerospace industry. The development of methodologies that can detect components, specially the smallest proportion of the rubbers composition is a constant challenge and an important gap in the studies of this nature. Therefore, methodologies by using standard techniques and/or of last generation are important in rubber industry and research laboratories, aiming application in related research. In this context, this study shows the coupling/association techniques (off-line) of thin layer chromatography and infrared spectroscopy (TLC/IR), being the IR spectra obtained by universal attenuated total reflection (UATR), applied to the analysis of additives in rubber compositions of ethylene-propylene-diene rubber (EPDM). Two EPDM compositions, a kind of eluent system and Gibbs’ reagent, as developer, were used. Basically, all organic components were detected by this methodology, being possible to suggest that it can be applied for detecting additives of similar chemical structures, even though it’s presents in small amounts in the composition. Keywords: additives, EPDM, TLC/IR off-line coupling, UATR.

1. Introduction The determination of the additive in polymers compositions in general by instrumental techniques, undergoes a preliminary extraction step. However, unambiguous identification of these components is not always possible, when analyzed by spectroscopic techniques, such as infrared spectroscopy (IR), due to the interference of other additives bands[1]. The research of different industrial goals requires the development of compositions containing plasticizers, accelerators and other additives, which may also form secondary products, increasing the complexity of the determination. The knowledge of the origin of the polymeric material, its application and suitable composition, are important data for the development of analytical methodology, for the characterization of these additives. A separation technique is needed to compose this methodology in the search for greater accuracy in obtaining results[2]. Thin layer chromatography (TLC) is a well known technique for separating components of a mixture and characterize unknown materials by comparison of the characteristics (color and reference- factor RF) of its separated products, related to reference materials, which must be shown in the composition[3]. TLC has advantages such as: speed, simplicity and low cost, it can be used in different laboratories that awakes greater interest of the industry and can be a useful tool for the scientific community. However, it is necessary to find suitable conditions for determining, which is not always a simple task[2,3].

74

IR spectroscopy, a well known technique[4], whose greatest quality is the identification of chemical function compounds, has been used quite successfully by many researchers, and in the laboratories of Instituto de Aeronáutica e Espaço (IAE), in particular, for analysis of rubber based on copolymer of ethylene-propylene-diene (EPDM), used in aerospace researches[5-7]. Regarding the analysis data of additives in rubber, papers have also been published by the IAE group, and among them some stand out, the use of liquid chromatography high performance (HPLC) and IR spectroscopy applied to study of rubber additives[8] in known composition on polychlorosulfoethylene (Hypalon) (CSM), and butadiene‑styrene copolymer (SBR) of unknown composition, and the application of technical coupling association of TLC/IR[2] for the study of rubbers. In the study[8], HPLC was used as a components separation technique, being confirmed by IR analysis of the CSM composition, the presence of tetramethylthiuram disulfide (TMTD). Mercaptobenzotialzol disulfide (MBTS) was not detected by IR, according to the authors, probably due to this compound is in lower content in the composition of rubber, therefore it was possible to confirm the presence of additives by HPLC with ultraviolet detection (UV). For SBR rubber, could be detected by IR, the presence of amine antioxidant, based on likely to diphenylamine and naphthenic oil. The presence of amine antioxidant, based on the similiarity chemical structure to diphenylamine and naphthenic oil, could be detected by IR in the SBR rubber.

Polímeros, 26(1), 74-80, 2016


TLC/IR (UATR) off-line coupling for the characterization of additives in EPDM rubber compositions HPLC analysis was not effective in this case. Thus, these techniques were complementary, in order to identify the greatest number of additives in the composition. Although some details of the compositions studied have been raised by this study[8], it is understood that complex method was used according to the conditions available, because HPLC analysis required several injections resulting in considerable time analysis. Dutra[2] described the TLC-IR coupling techniques as useful for the separation and identification of additives and uses two types of methods, one of pyramids including potassium bromide (KBr) and other simpler but also using KBr, developed in the laboratories of the IAE. It was observed that the application of this coupling or combination of techniques was made only using IR transmission techniques, with laboratory resources available during the study period. But, we think that is possible to include in to this study the use of reflection techniques. Recently, another study[5], Fourier transform infrared spectroscopy (FT-IR) was employed to investigate the gaseous products of EPDM. The potential of FT-IR analysis of gaseous pyrolyzates, named PY-G/FT-IR, for characterization of EPDM additives, was evaluated. Two EPDM compositions were analyzed. Initially, gaseous pyrolysis products from paraffinic oil, stearic acid, 2,2,4-trimethyl1,2-dihydroquinoline, tetramethylthiuram monosulfide (TMTM), TMTD, and 2-mercaptobenzothiazole (MBT) were characterized separately, and their main IR bands were identified. Subsequently, the gaseous pyrolysis products of three kinds of EPDM compositions, raw, unvulcanized, and vulcanized, were analyzed. The similarities observed in the FT-IR spectra of unvulcanized and vulcanized EPDM show that the vulcanization process does not interfere with the pyrolysis products. The identification of the functional groups of the studied additives was possible in both unvulcanized and vulcanized EPDM samples, without solvent extraction. It was possible to demonstrate that the PY-G/FT-IR technique can identify additives containing sulfur in concentrations as low as 1.4 phr in both unvulcanized and vulcanized EPDM. However, the methodology showed some limitations due to overlapping and similarities of TMTM and TMTD PY-G/FT‑IR spectra, which could not be distinguished from each other. Sanches et al.[5] considered that the PY-G/FT-IR is a faster and cheaper alternative technique compared to the sophisticated techniques usually applied to detection of additives in rubbers. However other techniques, such as TLC-IR, also can be applied to reach this goal. TLC association with the reflection techniques, used in Fourier transform spectroscopy (FT-IR), has been cited in the literature[9,10]. In these studies, by using reflection techniques, although it has been highlighted the diffuse reflectance (DRIFT), it was mentioned that there were difficulties related to time of analysis and interference of absorptions of materials used in DRIFT-TLC analysis, such as silica present on the TLC plate, with those relating to additives to be identified. Recently, FT-IR microscopy technique[11], have been focused as more promising, but we think that other reflection techniques, less complex and quicker can be used, such as universal attenuated total Polímeros, 26(1), 74-80, 2016

reflection (UATR), which has been used successfully to polymer analysis at IAE laboratories[6]. Therefore, in this paper, it was evaluated the applicability of the coupling/association TLC /IR techniques for the analysis of two EPDM rubber compositions, aiming to contribute to increase for increasing the knowledge for characterization of additives in rubber by potential reflection techniques, such as UATR.

2. Materials and Methods The samples used in the compositions of EPDM-were kindly supplied by the company Zanaflex Borrachas. In Table 1, the components of composition and the quantities (phr) are included, of the EPDM rubber compositions[12], numbers 1 and 2, with small content of components for internal purposes. After the development of the composition, samples were cured[10] with 2 mm thick plates per 10 cm of width and length, during 30 minutes at 160 °C.

2.1 Methodology / extraction The initial separation of organic additives EPDM of the two compositions was done by using auxiliary techniques, such as Soxhlet extraction with suitable solvent, acetone[13]. The material, about 0.5g, to be extracted was cut into small pieces and wrapped in filter paper and inserted to the extracting system. The time used for the extraction was 8 hours, and the solvent was evaporated under mild conditions.

2.2 TLC/IR analysis In an attempt to separate the organic additives, soluble in acetone, TLC technique has been used with Merck chromatography plate, 20 cm and it recovered with a layer of silica gel 60. The plate was marked in 0.78 inches for the upper and 0.39 cm for the lower, to limit the amount of eluent to be added in the glass tank[2]. In the chromatography plate was added 15 μL of the extract obtained[2] and the additives, taken as a reference: stearic acid, TMTD, MBT,TMTM, TMQ and paraffinic oil, being necessary to solubilize the Table 1. EPDM rubber compositions (1 and 2).

Componentes EPDM (Keltan 2340A) EPDM (Keltan 21) Zinc oxide Sulfur Stearic acid Carbon black Paraffinic oil MBT (mercaptobenzothiazole) TMTD (Tetramethyl thiuram disulfide)

TMQ (2,2,4- trimethyl-1,2dihydroquinoline) TMTM (Tetramethyl thiuram monosulfide)

Composition Composition 1 2 (phr) (phr) 100 100 5.0 2.0 1.5 0.7 1.0 0.5 80 5.0 50 1.0 0.5 0.7 1.0 --

1.0

-

0.7

75


Damazio, D., Campos, E. A., Diniz, M. F., Mattos, E. C., & Dutra, R. C. L. starting materials, in powder form, with a suitable solvent for the formation of liquid. Some analyzes were performed in duplicate to verify system repeatability. After labeling and addition of materials, the plate was placed in the chromatography vessel saturated with the chosen eluent system, hexane, diethyl ether and acetic acid[14] (70: 30: 5), in the 1 cm mark to occur the development of analysis[2]. After the run time, suitable for each system analyzed, the plate was removed and brought to 105 °C for 30 minutes in an oven, and then cooled to room temperature. After the drying and cooling time, the plate was subjected to development with a suitable developer spots of 1% solution of 2,6-di-chloro-para-chloro benzoquinine imine (Gibbs reagent)[2,12], which acts revealing the colors, highlighting the differences of the analyzed additives. Each additive was eluted by a certain height. For all the products analyzed, the ratio was measured height, in inches, of the eluted deposit average and total height of 16 cm. This value is called retention factor (RF). The eluted deposit obtained on the plate was scraped and treated with the solvent used in the extraction, filtered and dried to remove silica[2] and then analyzed by using FT-IR spectrometer SPECTRUM ONE PERKINELMER. The analysis conditions were: resolution: 4 cm–1; in the mid infrared (MIR), 4000 to 400 cm–1; 20 scans, using the reflection accessory, UATR. The additives were also analyzed as received by UATR under the same conditions.

3. Results and Discussions 3.1 UATR analysis of the organic additives as received Before proceeding to step methodology involving the extraction and characterization of the components by TLC / FT-IR, it is convenient to perform, initially, the characterization of the reference additives according to the same method of obtaining spectra, by UATR. Figure 1 shows the UATR spectra of composition additives and their structural formulas.

according to the same method, it constitutes the second step of the developed methodology. The major bands observed (cm–1) in UATR spectrum of the extract in acetone of composition 1 shows the presence of the functional groups[4]: OH (3400 cm–1), CH2 (2922 and 2853), C=O (1713), CH2 and CH3 (1457), CH3 (1377), C-O and/or C-S (1243), C-O (1164 and 1099) and CH2 n ≥3 (720), suggesting that the residue consists essentially of aliphatic compound containing acid groups (COOH), CO or CS, and carbon chain with more than 3 CH2 groups in accordance with some compounds present in the composition, such as stearic acid, TMTD, but, as expected, may there are overlapped bands, requiring a separation technique such as TLC, and other identification step, such as FT-IR technique, in trying to get a better result. The major bands observed in UATR spectrum of the extract in acetone, composition 2, contains the wavenumbers (cm–1)[4]: 3373 (OH or NH), 2959 (CH3), 2924 and 2854 (CH2), 1709 (C=O), 1605 and 1498 (ArC-C and CHN), 1461 (CH2 and CH3), 1379 e 1360 (CH3 gem-dimethyl), 1319 (C-N and/or C=N), 1245 (C-O and/or C-S, C-N C=S), 1103 (C-O), 1015 (C-O and/or ArC-S), 815, 750 and 655 (ArC-H). This set of bands suggests that the residue of the composition 2, consisting primarily of aliphatic and aromatic compounds containing NH groups and CS, similar to those found in sulfur and amino compounds in the composition, such as MBT, TMQ and TMTM and also the presence of acid groups COOH and CO. However, in the same manner as for the composition 1, may overlap bands, requiring the application of a separation technique such as TLC and an identification technique, such as FT-IR.

The main bands, assigned to its main functional groups[4,15] are in (cm–1): a) MBT - 3108, 3070 and 3037 (ArCH), 1595 and 1495 (ArCC), 1423 (C=N and/or C-S), 1319 and 1243 (C=N, C-N and /or C-S), 1012 (ArC-S), 749 (ArCH); b) paraffinic oil - 2921 (CH2), 2852 (CH2), 1460 (CH3, CH2), 1376 (CH3) and 720 (CH2 n ≥3); c) stearic acid – 1698 (COOH), 1296 (CO), 720 (CH2 n ≥3); d) TMTD - 2933 (CH3), 1370 (CH3), 1232 (C=S, C-S), 1145 (C-N), 563 (SS); e) TMQ - 3380 (NH), 3017 (ArCH), 2961 (CH3), 2926 e 2866 (ArCH3), 1649 NH, 1607 (ArCH), 1581 (NH), 1499 (ArCH, CHN), 1381 and 1361 (CH3 gem-dimethyl), 1260 and 1170 (ArC-N), 814 and 748 (ArC-H); f) TMTM - 1518 (C=S, C-N), 1438, 1249, 1149, 1052 e 998 (C-S, C=S, C-N),1376 (CH3).

3.2 UATR analysis of extraction residue of the two EPDM compositions Since the method chosen for characterizing eluted spots on TLC/IR method was UATR, the extraction residues of EPDM rubber compositions 1 and 2 (Figure 2) were analyzed 76

Figure 1. UATR spectra of organic additives of EPDM compositions 1 and 2: (A) MBT; (B) paraffinic oil; (C) stearic acid; (D) TMTD; (E) TMQ; (F) TMTM. Polímeros, 26(1), 74-80, 2016


TLC/IR (UATR) off-line coupling for the characterization of additives in EPDM rubber compositions

Figure 2. Spectra UATR of extraction residue in acetone, of EPDM compositions: (A) 1; (B) 2.

It should be recorded that the spectra of Figure 2 is different and may already be a sign that are associated to different compositions, constituting a first information and support to analysis, if they were analyzing unknown compositions, for example.

3.3 TLC analysis of EPDM compositions 3.3.1 Composition 1 - TLC system: eluent containing 70% hexane, 30% ethyl ether and 5% acetic acid / Gibbs developer (Bukhina et al.[14]) Figure 3 shows the chromatography plate of composition 1, after using the TLC system mentioned. The RFs and colors of spots, shown in Figure 3 already mentioned, are included in Table 2. Usually referred call the first eluted deposit of each sample, 1, and to give sequential numbers to others, for this and the other composition .Although there is no very clear separation of colors, very clear, can be seen that the sample extract showed basically an eluted deposit. Additives showed two eluted spots, each showing good repeatability in color and RF. The stearic acid showed one eluted deposit. There is an indication, the RF values and colors, that both paraffinic oil (higher RF) and MBT (RF around 0.4) would be present in the sample extract, but there is need for the IR analysis to try to confirm the indication TLC analysis. Polímeros, 26(1), 74-80, 2016

Figure 3. Composition 1 - TLC system: mixture of 70% hexane, 30% ethyl ether and 5% acetic acid in a ratio of 70:30:5 and Gibbs reagent as developer.

3.3.2 Composition 2 - TLC system: eluent containing 70% hexane, 30% ethyl ether and 5% acetic acid / Gibbs developer (Bukhina et al.[14]) Figure 4 shows the chromatography plate of composition 2, after using the TLC system above named. The RFs and colors of spots, shown in Figure 4 cited above, are included in Table 3. It is observed that there was 77


Damazio, D., Campos, E. A., Diniz, M. F., Mattos, E. C., & Dutra, R. C. L. Table 2. TLC system – Composition 1: eluent of 70% hexane, 30% ethyl ether and 5% acetic acid), after using Gibbs’ reagent - spots of the composition extract and their additives. Nº. Eluted dep.

RF calculated

Colour

A - acetone extract

1.0

1.00

light red

B - acetone extract

1.0

1.00

light red

acetone extract middle of the plate

1.0

0.31

light brown

C - MBT

1.0

0.37

dark brown

2.0

0.47

light brown

1.0

0.25

dark brown

2.0

0.47

light brown

1.0

0.14

light brown

2.0

0.25

dark brown

1.0

0.14

light brown

2.0

0.25

dark brown

G - Stearic acid

1.0

0.65

dark beige

H - Paraffinic oil

1.0

1.00

light beige

Sample

D - MBT E - TMTD F - TMTD

Table 3. TLC system- Composition 2: mixture of 70% hexane, 30% ethyl ether and 5% acetic acid in a ratio of 70:30:5 and Gibbs reagent as developer - spots of the composition extract and their additives. Sample A - Acetone extract

B - Stearic acid C - Paraffin oil D - MBT E - TMQ

F - TMTM

N0. eluted dep. 1 2 3 4 5 6 1 1 1

RF calculated 0.03 0.2 0.25 0.38 0.44 0.69 0.53 0.78 0.2

1 2 3 1

0.1 0.44 0.69 0.11

Colour Blue Blue Brown Brown Blue Blue Light beige Light brown Reddish Brown Blue Blue Blue Yellow

a clearer separation, color and RF. In this plate, in addition to the additives already discussed in another composition and that are common to both, there was also the addition of TMQ and TMTM, components of the composition 2. There is an indication by the RF values and colors, as well MBT would TMQ present in the sample extract, but there is also need for the IR analysis, as mentioned for other formulation to try to confirm the indication of TLC analysis.

3.4 TLC/IR off-line coupling of EPDM compositions TLC/IR[2] methodology consists basically in assessing the existence of spots of rubber extract with colors and RF similar to those of organic additives, the chromatographic 78

Figure 4. Composition 2 - TLC system: mixture of 70% hexane, 30% ethyl ether and 5% acetic acid in a ratio of 70:30:5), and Gibbs reagent as developer.

plate studied, shave them, trying to eliminate interference of the board of silica and obtaining IR spectra of these spots, effecting thereby a comparison spectra, in order to characterize the additives present in the formulation. Therefore, the spectra shown below are for the comparison made for both EPDM formulations. It must be remembered that the characterization of inorganic additives, if necessary, is made separately for simple calcination[2]. • Composition 1: Figure 5 shows the UATR spectra of spots of extracts the A and B rubber, compared to others of the TMTD, stearic acid, paraffin oil and MBT, also suggested by UATR analysis of the rubber extract or by TLC analysis. In this figure and in the other composition, were marked only the bands that distinguish the functional groups of the compounds, preventing common absorptions noted CH2 and CH3 groups, and some that are in regions close to the CS as CN, around 1150 cm–1, in order to facilitate the visualization of bands and explanation of the data. However, it can be noted that the spectrum of the extracts A and B - middle of the plate, shows similarity with the shape of absorptions of sulfur compounds, even though in a smaller proportion in the formulation.

Basically, it can be said that were observed bands (cm–1) of stearic acid, about 1700 (C=O) and 1300 (CO), MBT, around 3070 to 1600 (Ar-H and Ar-C); TMTD of around 1240 (C=S, CS), and 560 (SS); of paraffin oil, around 720, confirming the information of UATR and TLC analysis of extract of this composition. • Composition 2: Spectra were obtained from different RF composition extract, but some were not well resolved probably due to the small quantity of sample. There were in others interference of silica, and basically, additives which might be characterized were stearic acid (bands around 1700 cm–1 (C = O) and 1300 cm–1 (CO), and TMQ (bands around 3380 cm–1 (NH) and 741 cm–1, (ArC) as shown in Figure 6. Polímeros, 26(1), 74-80, 2016


TLC/IR (UATR) off-line coupling for the characterization of additives in EPDM rubber compositions

Figure 5. UATR Spectra of spots of the composition 1: (A) A RF1; (B) B, RF1; (C) A and B - middle of the plate; (D) stearic acid; (E) MBT; (F) TMTD; (G) paraffinic oil.

Figure 6. UATR spectra of spots of the composition 2): A and B, RF5; (B) stearic acid; (C) A and B, RF2; (D) TMQ. Polímeros, 26(1), 74-80, 2016

79


Damazio, D., Campos, E. A., Diniz, M. F., Mattos, E. C., & Dutra, R. C. L.

4. Conclusions TLC/IR (UATR) off-line coupling showed that it is possible to separate and identify some additives, even in small quantities in the EPDM compositions analysed, by a simple and fast IR methodology, without interference of silica IR bands, as observed by the others more complex techniques such as HPLC/IR or TLC/IR (DRIFT or Microscopy/FT-IR). Thus, it is possible to suggest that this methodology could be applied for detection of organic additives of similar chemical structures, even in small quantities present in the rubber composition.

5. References 1. Crompton, R., & Crompton, T. R. (2007). Determination of additives in polymers and rubbers. Southport: Smithers Rapra Publishing. 2. Dutra, R. C. L. (1996). Aplicação de técnica TLC-IR em estudos de separação, identificação e quantificação de aditivos em borrachas. Polímeros: Ciência e Tecnologia, 6(2), 26-31. Retrieved in 13 February 2015, from http://revistapolimeros. org.br/files/v6n2/v6n2a01.pdf 3. Striegel, M. F., & Hill, J. (1996). Thin-layer chromatography for binding media analysis (Scientific Tools for Conservation). Los Angeles: Getty Conservation Institute. Retrieved in 13 February 2015, from http://hdl.handle.net/10020/gci_pubs/ thin_layer_chromatography 4. Smith, A. L. (1979). Applied infrared spectroscopy. New York: John Wiley & Sons. 5. Sanches, N. B., Cassu, S. N., Diniz, M. F., Dutra, R. C. L. (2014). Characterization of additives typically employed in EPDM formulations by using FT-IR of gaseous pyrolyzates. Polímeros: Ciência e Tecnologia, 24(3), 269-275. http://dx.doi. org/10.4322/polimeros.2014.066. 6. Santos, R. P., Oliveira, M. S., Mattos, E. C., Diniz, M. F., & Dutra, R. C. L. (2013). Study by FT-IR technique and adhesive properties of vulcanized EPDM modified with plasma. Journal of Aerospace Technology and Management, 5(1), 65-74. http:// dx.doi.org/10.5028/jatm.v5i1.162. 7. Santos, R. P., Oliveira, M. S. Mattos, E. C.; Diniz, M. F.; Dutra, R. C. L. (2012). Caracterização por FT-IR da superfície de

80

borracha EPDM tratada via plasma por micro-ondas. Polímeros: Ciência e Tecnologia, 22(5), 440-446. http://dx.doi.org/10.1590/ S0104-14282012005000065. 8. Dutra, R. C. L., & Montenegro, A. M. C. (1987). Análise de aditivos em borrachas. In: Anais do I Seminário de Caracterização de Sistemas Poliméricos. Rio de Janeiro. 9. He, W., Shanks, R., & Amarasinghe, G. (2002). Analysis of additives in polymers by thin-layer chromatography coupled with Fourier transform-infrared microscopy. Vibrational Spectroscopy, 30(2), 147-156. http://dx.doi.org/10.1016/ S0924-2031(02)00024-3. 10. Cserháti, T., Forgács, E., Candeias, M., Vilas-Boas, L., Bronze, R., & Spranger, I. (2000). Separation and tentative identification of the main pigment fraction of raisins by thinlayer chromatography-fourier transform infrared and highperformance liquid chromatography-ultraviolet detection. Journal of Chromatographic Science, 38(4), 145-150. http:// dx.doi.org/10.1093/chromsci/38.4.145. PMid:10766480. 11. Liu, J., Zhou, F., Guo, R., Jiang, Y., Fan, X., He, A., Zhai,Y., Weng, S., Yang, Z., Xu,Y., Noda, I., Wu, J. (2014). Analysis of an alanine/arginine mixture by using TLC/FTIR technique. Journal of Spectroscopy, 2014, 1-4. http://dx.doi. org/10.1155/2014/925705. 12. American Society for Testing and Materials – ASTM. (2013). ASTM D3568-03: standard test methods for rubber: evaluation of EPDM (Ethylene Propylene Diene Terpolymers) including mixtures with oil. West Conshohocken. Reapproved 2009. 13. Wake, W. C., Tidd, B. K., & Loadman, M. J. R. (1983). Analysis of rubber and rubber-like polymer. 3rd ed. New York: Applied Science. 14. Bukhina, M. F., Morozov, Y. L., van de Ven, P. M., & Noordermeer, J. W. M. (2003). Mould fouling of EPDM rubber compounds. KGK Kautschuk Gummi Kunststoffe, 56(4), 172-183. Retrieved in 13 February 2015, from: http://www.kgk-rubberpoint.de/ ai/resources/e907ea3cdb1.pdf 15. Wolfang, W. (1987). Tópicos de espectroscopia no infravermelho. São José dos Campos: ITA. Received: Feb. 13, 2015 Revised: July 20, 2015 Accepted: Oct. 01, 2015

Polímeros, 26(1), 74-80, 2016


http://dx.doi.org/10.1590/0104-1428.2087

Otimização do processo de dispersão de nanotubos de carbono em poliuretano termorrígido Optimization of carbon nanotubes dispersion process in thermoset polyurethane Magnovaldo Carvalho Lopes1, João Paulo Campos Trigueiro1, Vinicius Gomide de Castro1, Rodrigo Lassarote Lavall1 and Glaura Goulart Silva1* Laboratório de Materiais Poliméricos Multicomponentes, Departamento de Química, Universidade Federal de Minas Gerais – UFMG, Minas Gerais, MG, Brasil

1

glaura.goulart@gmail.com

Resumo Neste trabalho foi desenvolvido um processo empregando misturador de alto cisalhamento e moinho de rolos para dispersar MWCNTs (multiwalled carbon nanotubes) puros e modificados em poliol visando a preparação de concentrados de 3% em massa. Condições otimizadas no trabalho permitiram a obtenção de suspensões com menor número e tamanho de agregados de MWCNTs. Compósitos contendo 0,5% em massa de MWCNTs foram preparados por diluição dos concentrados em poliol usando mistura mecânica seguida de cura. Resultados de microscopia indicaram que as melhores dispersões foram obtidas com os MWCNTs modificados, os quais permitiram um aumento na tensão na ruptura, no alongamento e uma melhor preservação da estabilidade térmica. Além disso, valores de condutividade elétrica sugerem que o compósito possa ser empregado para dissipação eletrostática. Dessa forma, os resultados obtidos demonstram que a modificação covalente da superfície dos MWCNTs e a utilização de estratégias eficientes de dispersão são essenciais para melhorar as propriedades finais dos nanocompósitos. Palavras-chave: dispersão de nanotubos de carbono, nanotubos de carbono modificados, propriedades mecânicas, poliuretano termorrígido elastomérico. Abstract A process employing high shear mixer and roll mill to disperse pristine and modified MWCNTs (multiwalled carbon nanotubes) in polyol was developed in order to prepare 3 wt% masterbatches. The optimum process conditions resulted in suspensions with smaller number and size of nanotube aggregates. Composites containing 0.5 wt% of MWCNTs were prepared by dilution of polyol masterbatches by simple mechanical mixing followed by cure. Microscopy data revealed better dispersion of modified carbon nanotubes in the polymer matrix, which promoted an increase in the tensile strength, elongation and a better preservation of thermal stability. Furthermore, electric conductivity values indicated that the composites can be used for electrostatic dissipation. These results demonstrate that the covalent modification of MWCNTs surface and the use of efficient dispersion strategies are essential to improve nanocomposites’ final properties. Keywords: carbon nanotubes dispersion, modified carbon nanotubes, thermoset polyurethane elastomer, mechanical properties.

1. Introdução A primeira descrição relacionada ao processo de produção de poliuretanos (PUs) baseado na reação de um diisocianato com um diol foi a patente alemã depositada pela I.G. Farben (subdivisão da Bayer) em 1937, tendo Otto Bayer como um dos inventores.[1] As propriedades dos PUs e portanto sua aplicação final, dependem do isocianato e diol empregados (e da proporção entre eles), além de outros reagentes como extensores de cadeia, agentes de cura, aditivos, entre outros, bem como do processo envolvido na sua síntese. Esse processo pode acontecer em uma etapa (one shot), no qual todos os

Polímeros, 26(1), 81-91, 2016

reagentes são adicionados no início da reação ou em duas etapas, via produção de um pré-polímero (reação entre o isocianato e poliol) que é posteriormente reagido com o extensor de cadeia ou agente de cura (sistemas com um ou dois componentes)[1,2]. Podem ser produzidos poliuretanos termoplásticos ou termorrígidos, preparando-se desde espumas rígidas ou flexíveis, passando por revestimentos resistentes a produtos químicos, adesivos especiais, selantes, pequenos componentes de máquinas como engrenagens, rodízios e roldanas, até grandes peças industriais de alto desempenho mecânico[1,2,3].

81

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


Lopes, M. C., Trigueiro, J. P. C., Castro, V. G., Lavall, R. L., & Silva, G. G. A presença de ligações covalentes entre as cadeias (ligações cruzadas) fornecem aos PUs termorrígidos propriedades superiores às verificadas para os PUs termoplásticos como resistências à tração, compressão, impacto, abrasão e à degradação causada por ácidos, bases e solventes orgânicos, mantendo as características elastoméricas[4]. Esses polímeros são geralmente preparados via processamento de pré-polímero, uma vez que possibilita um melhor controle da reação química antes da etapa de reticulação, empregando-se proporções molares da ordem de 1,00 mol de diol para 1,60-2,25 mols de isocianato[2,3]. Os poliuretanos termorrígidos elastoméricos com excelentes propriedades de resistência ao desgaste são utilizados nas indústrias de mineração, óleo e gás, apresentando grandes vantagens financeiras, sendo utilizados na fabricação de peças estruturais de alto desempenho tais como esteiras de mineração, laminados de interior de aviões ou peças de recepção de grandes tubulações em alto mar (enrijecedores de curvatura)[5,6]. Os nanotubos de carbono (CNTs) são considerados como um dos materiais mais resistentes já sintetizados pelo homem. Apresentam propriedades elétricas e térmicas extraordinárias somadas a uma boa estabilidade química, o que vem despertando grande interesse na sua utilização em diversas aplicações, principalmente no preparo de compósitos poliméricos com elevada resistência mecânica, por exemplo à tração e flexão, assim como grande estabilidade térmica[7,8]. Os nanotubos de carbono de paredes múltiplas (MWCNTs) possuem características elétricas entre metal e semicondutor, apresentam módulo de elasticidade de 0,27‑0,95 TPa, resistência à tração de 11-63 GPa e condutividade entre 200-3000 W/mK (MWCNTs isolados)[9-11]. A adição de CNTs em matrizes poliméricas permite melhorias nas suas propriedades mecânicas, térmicas, elétricas, entre outras[10-12]. A redução no preço dos nanotubos de carbono, especialmente para os MWCNTs, tem contribuído para sua utilização em maiores escalas[13]. Como o processamento dos poliuretanos termorrígidos é completamente distinto, a dispersão de CNTs nessa matriz para o preparo de compósitos exige técnicas de misturas específicas e seu estudo ainda encontra-se em fase incipiente quando comparado às pesquisas com PUs termoplásticos. Além das características relacionadas ao nanotubo e ao poliuretano discutidas anteriormente, o processo empregado na dispersão dos CNTs é determinante para as propriedades finais do compósito obtido. Os poucos trabalhos voltados para o preparo de nanocompósitos CNTs/PUs termorrígidos encontrados na literatura são baseados na síntese do polímero em duas etapas (via pré-polímero) com a dispersão dos nanotubos no isocianato ou no poliol com o auxílio de dispersores de alto cisalhamento ou ultrassom[14-17]. Foram utilizados nanotubos do tipo SWCNTs (nanotubos de carbono de parede simples) e MWCNTs, purificados, modificados e não modificados quimicamente[15-18]. Nesses trabalhos, os compósitos foram preparados diretamente na concentração de interesse. No intuito de contribuir para o tema e acrescentar novos dados ao universo de compósitos com poliuretanos termorrígidos, nosso grupo propôs uma metodologia de preparação de compósito PU termorrígido elastomérico/MWCNT por meio 82

da dispersão dos MWCNTs em pré-polímero empregando dispersor de alto cisalhamento e moinho de rolos a partir de concentrados (masterbatches) que posteriormente foram diluídos para a concentração desejada[19,20]. Como continuação desse estudo, o presente trabalho tem como objetivo avaliar um procedimento de dispersão de CNTs por meio da avaliação da morfologia, propriedades térmicas, elétricas e mecânicas dos nanocompósitos obtidos. Foram utilizados MWCNTs não modificados e modificados (com grupos carboxílicos) em poliol por meio de preparação de concentrados (3,0% em massa de MWCNTs). Esse trabalho visa contribuir para a introdução de nanotubos de carbono no setor produtivo de poliuretanos termorrígidos buscando diminuir o impacto negativo (financeiro e de saúde, meio ambiente e segurança - SMS) na cadeia de valores do produto final.

2. Materiais e Métodos 2.1 Materiais Foram utilizados MWCNTs adquiridos da empresa Timesnano (China) com pureza >95%, diâmetro externo e comprimento de 2-8 nm e 1-30 μm, respectivamente. Dois tipos de nanotubos foram utilizados: não modificado e modificado com grupos carboxilados (4% de acordo com o fornecedor). Para a síntese do PUE (poliuretano termorrígido elastomérico), foram empregados poli (tetrametileno éter glicol) (PTMG-1000) proveniente da SAFE Chemicals LLC, 2,4-diisocianato de tolueno (TDI) (nome IUPAC 2,4-diisocianato -1-metil benzeno) fornecido pela Bayer S.A., 1,4-butano diol como extensor de cadeia (nome IUPAC butano-1,4-diol) proveniente da M.Cassab Comercio e Indústria LTDA e 4,4-metileno-bis-orto-cloroanilina (MOCA) como agente de cura (nome IUPAC [(4-amino‑3‑clorofenil) metil]-2-cloroanilina) adquirido da M.Cassab Comercio e Indústria LTDA. Todos os reagentes foram utilizados como recebidos.

2.2 Métodos 2.2.1 Dispersão de MWCNTs em poliol Inicialmente os MWCNTs modificados e não modificados foram secos em estufa a vácuo por um período de 10 horas a temperatura de 110 °C de forma a garantir a ausência de umidade. Posteriormente, a massa desejada de nanotubos (modificados e não modificados) foi dispersa em PTMG‑1000 previamente aquecido a 60 °C. Esta mistura originou um concentrado (masterbatch) com teor de 3% em massa de MWCNTs em Poliol. A dispersão foi realizada através de um misturador de alto cisalhamento (Turrax) a uma velocidade de 20000 rpm por 10 min. Após esta etapa a mistura foi processada em moinho de rolos com separação entre os rolos de 10 e 5 µm e velocidade de 100 rpm a 60 °C. A Figura 1 mostra as etapas usadas na dispersão. 2.2.2 Síntese dos nanocompósitos Os nanocompósitos foram produzidos em um processo que envolve a síntese de um pré-polímero e sua posterior cura (Figura 1b). A reação para obtenção do pré-polímero foi realizada em um reator de aço inox equipado com agitador mecânico sob vácuo após diluição do concentrado (pela adição de PTMG) para 0,5% em massa de MWCNTs. Foram adicionadas Polímeros, 26(1), 81-91, 2016


Otimização do processo de dispersão de nanotubos de carbono em poliuretano termorrígido

Figura 1. (a) Processo de dispersão dos nanotubos de carbono em poliol em duas etapas, primeiro com o agitador mecânico e posteriormente com o moinho de rolos e (b) produção dos compósitos[21].

quantidades estequiométricas de PTMG 1000/MWCNTs/TDI e o sistema foi mantido a 70 °C por 2 horas com posterior adição de 1,4-butanodiol com a finalidade de aumentar a massa molar do pré-polímero[22]. Após 2 horas de reação, obteve-se o pré-polímero com teor de NCO livre em torno de 7,4% determinado por titulação com N-dibutilamina[23]. A segunda etapa do processo de síntese dos nanocompósitos inicia-se com a desaeração do pré-polímero em câmara de vácuo por 90 min. Antes de adicionar o agente de cura (MOCA) o pré-polímero foi reaquecido a 80 °C. A massa de MOCA utilizada foi proporcional à quantidade de NCO livre presente no pré-polímero. Após a adição do MOCA, a mistura foi homogeneizada através de uma agitação suave por aproximadamente 1 min tomando-se o devido cuidado de não inserir bolhas no sistema. Em seguida a mistura foi vertida em moldes que, posteriormente, passaram pelo processo de cura a 100 °C por 10 h. Posteriormente o processo de pós-cura final foi conduzido mantendo as amostras por 15 dias a temperatura ambiente[20]. 2.2.3 Caracterização A caracterização detalhada das amostras de nanotubos modificados e não modificados foi descrita em um trabalho prévio do nosso grupo[20]. Polímeros, 26(1), 81-91, 2016

Os nanocompósitos foram caracterizados por microscopia óptica (MO) utilizando um microscópio Olympus, modelo BX50F-e. Sua morfologia foi examinada por microscopia eletrônica de varredura (MEV) utilizando microscópio com canhão de emissão por efeito de campo (Quanta 200 - FEG/FEI). Para esta análise as amostras foram fraturadas em nitrogênio líquido e suas superfícies de fratura foram recobertas por uma fina camada de ouro. Com o propósito de avaliar a dispersão micrométrica e nanométrica dos MWCNTs na matriz de PUE foi utilizada microscopia eletrônica de transmissão (MET) em microscópio Tecnai – G2-20/FEI. Para esta análise, as amostras foram preparadas com o auxílio da técnica de criomicrotomia. Espectros de absorção no infravermelho do PUE e dos nanocompósitos foram obtidos utilizando espectrômetro FTIR (dispositivo de ATR, modelo 380 Nicolet Thermo Scientific) com uma faixa espectral de 4000-680 cm–1. Foram realizadas 32 varreduras com resolução de 4 cm–1. Medidas elétricas nestas mesmas amostras foram feitas a 25 °C utilizando um potenciostato Autolab PGSTAT30 Ecochemie. Foi utilizada uma célula experimental de impedancimetria modelo AN8080 da marca Analógica possuindo dois eletrodos bloqueantes de aço inox de área 0,27 cm2. A faixa de frequência usada foi entre 1 MHz e 0,5 Hz, sob 0 V e amplitude de perturbação 83


Lopes, M. C., Trigueiro, J. P. C., Castro, V. G., Lavall, R. L., & Silva, G. G. de 10 mV. As amostras foram medidas na forma de filmes com espessura de 0,3 mm. O PUE e os nanocompósitos foram submetidos a ensaios de tração e medição da dureza (Shore D). Os ensaios de tração foram realizados em máquina de ensaio Kratos modelo TRCv59D-USB com célula de carga de 100 Kgf, velocidade de deslocamento de 500 mm/min a temperatura de 22,0 °C e umidade relativa de 44,0% empregando-se um número mínimo de 03 ensaios para cada amostra. Análises termogravimétricas (TGA) destas amostras foram realizadas em atmosfera de nitrogênio com razão de aquecimento de 10 °C/min utilizando o equipamento TGA Q5000 dentro de um intervalo de temperatura de 20 a 800 °C.

3. Resultados e Discussões 3.1 Dispersão dos MWCNTs em poliol Na síntese de poliuretanos empregando-se reação em duas etapas, uma primeira avaliação da dispersão dos nanotubos de carbono pode ser realizada através de imagens

obtidas por microscopia ótica. Essa técnica é muito útil para verificar se há presença de agregados de CNTs, inferir sobre o número e tamanho dos mesmos e assim julgar a eficiência do processo de dispersão utilizado[24,25]. A Figura 2 mostra as etapas do processo de dispersão dos nanotubos de carbono obtidas no presente trabalho e as imagens de microscopia óptica do sistema em cada etapa. Como descrito na literatura, o grande desafio no preparo de nanocompósito poliméricos é a dispersão eficiente dos nanotubos de carbono na matriz, tanto em suspensão quanto no estado sólido[26]. A imagem de MO (Figura 2b) mostra a sedimentação dos nanotubos de carbono (MWCNT modificado) após sua adição ao poliol e agitação simples. Para diminuir a viscosidade do meio e melhorar a processabilidade, o sistema foi aquecido a 60 °C. Na primeira etapa, a suspensão foi submetida a um dispersor de alto cisalhamento (Turrax) e a melhor dispersão foi obtida após 10 min na velocidade de 20000 rpm. No entanto, constata-se que existem vários agregados de diferentes tamanhos (na faixa de 3 µm ate 67 µm) de nanotubos de carbono não dispersos, o que indica

Figura 2. Processo de dispersão dos nanotubos de carbono em poliol em duas etapas: primeiro com o agitador mecânico e posteriormente com o moinho de rolos. (a) Poliol e nanotubo formando duas fases, (b) imagem de MO da dispersão dos CNTs em poliol utilizando apenas agitador mecânico, (c) imagem de MO da dispersão dos CNTs em poliol após as duas etapas de uso do agitador mecânico e do moinho de rolos, (d) CNTs dispersos em poliol formando uma única fase. 84

Polímeros, 26(1), 81-91, 2016


Otimização do processo de dispersão de nanotubos de carbono em poliuretano termorrígido que apenas a utilização do dispersor não é suficiente para atingir o efeito desejado. Na sequência, o sistema foi levado a um moinho de rolos e submetido a diferentes condições de processamento, variando-se a separação entre os rolos, a velocidade de rotação do moinho e o número de vezes pelo qual determinada quantidade de material era submetida aos três rolos (“ciclos”). A imagem de MO (Figura 2c) mostra a dispersão obtida (MWCNT modificado) com separação entre os rolos de 10 e 5 µm, velocidade de rotação de 100 rpm após 10 ciclos. Embora ainda existam pequenos agregados (entre 0,7 e 21 µm), verifica-se que a quantidade de material disperso aumentou com a introdução da etapa de moagem. O aspecto da dispersão final é mostrado na Figura 2d. Para o MWCNT não modificado, esta também é a melhor condição de dispersão, embora tenham sido observados agregados maiores e em maior número que para o MWCNT modificado. A melhor dispersão dos MWCNTs modificados está relacionada à sua melhor interação com o PTMG, o que previne a reagregação. A vantagem do processo de dispersão do nanotubo em poliol em relação ao realizado em pré-polímero é que o controle da umidade não precisa ser rigoroso, o que facilita o transporte e armazenamento do produto. Como já mencionado, a proporção isocianato/poliol é crucial para a obtenção de PUE com propriedades desejadas. Após a reação de formação do pré-polímero, a manutenção da quantidade de isocianato livre calculada para permitir a reação com o agente de cura é importante. A presença de vapor de água altera consideravelmente o teor de NCO livre, e, portanto, exerce grande influência nas características do compósito final. As imagens de MO mostradas na Figura 2 são relativas ao concentrado (masterbatch) contendo 3% em massa de CNT funcionalizado. Esse concentrado pode ser armazenado por vários dias sem que nenhuma sedimentação ocorra. Para preparo dos compósitos, o masterbatch pode ser diluído para a concentração desejada pelo emprego de agitação mecânica simples. Isso facilita a introdução dos CNTs na cadeia produtiva de produtos/peças de poliuretanos termorrígidos, uma vez que não há necessidade de utilização de equipamentos específicos, pois os CNTs já estarão devidamente dispersos no concentrado, nem problemas relacionados à SMS, porque não haverá risco de presença de material particulado sólido suspenso no ar no ambiente de produção.

3.2 Caracterização dos nanocompósitos Para o preparo dos nanocompósitos, os concentrados contendo ambos os MWCNTs (modificados e não modificados) foram diluídos em poliol para concentração de 0,5% em massa e curados. Os materiais foram caracterizados após 15 dias de pós-cura (temperatura ambiente). A dispersão dos MWCNTs na matriz e a morfologia dos nanocompósitos foram avaliadas por MO, MEV e MET. Na Figura 3a, b são apresentadas as imagens de MO para os compósitos preparados com MWCNTs modificados e não modificados. Pode-se constatar que não houve mudança substancial na dispersão/distribuição dos nanotubos no estado sólido em relação aquela verificada para as suspensões em PTMG, ou seja, com as etapas de reação para a síntese de Polímeros, 26(1), 81-91, 2016

pré-polímero e cura não houve reagregação dos nanotubos de carbono. Tal fato é resultante da eficiência do processo empregado na dispersão para ambos os nanotubos. Como nas suspensões, pode-se observar a presença de aglomerados menores de MWCNTs modificados (entre 0,6 e 12 µm) em relação aos MWCNTs não modificados (faixa de 0,8 µm ate 27,0 µm), provavelmente devido a sua melhor interação com a matriz de PUE. As imagens de MEV (Figura 3c, d) corroboram o que foi pressuposto a partir das imagens de MO. Percebe-se que os MWCNTs modificados apresentaram uma melhor adesão à matriz polimérica como resultado da interação MWCNT-COOH:PUE. Além disso, como destacado por McClory e colaboradores[15], outro indício dessa interação é que as imagens foram obtidas na superfície de fratura e não foram observadas regiões que indicassem que o MWCNT foi puxado para fora da matriz no momento da fratura. De fato, essa interface MWCNT:matriz é essencial para o reforço mecânico no sistema. As imagens das Figura 3c, d exibem MWCNTs curvados e entrelaçados na matriz polimérica, ligando “pontos” da superfície fraturada. Essa morfologia, verificada em outros compósitos nanotubo:PUs (termorrígido ou termoplástico), é essencial para a transferência de tensão entre os CNTs e a matriz quando o material é sujeito a uma solicitação mecânica[15,27]. Como evidência do recobrimento pelo polímero, o diâmetro dos nanotubos de carbono visíveis nas imagens é da ordem de 23 nm a 34,0 nm nos compósitos com MWCNTs não modificados e de 12,0 nm a 36,0 nm nos compósitos com MWCNTs modificados. Valores bem menores foram observados para esses nanotubos isolados (diâmetro na faixa de 5,0 nm a 15,0 nm). Além disso, como os MWCNTs possuem apenas 4% de funcionalização, não deve haver prejuízo considerável às suas propriedades elétricas, e a morfologia verificada também é propícia à obtenção de boas propriedades elétricas para os compósitos. A melhor dispersão dos MWCNTs modificados também foi evidenciada nas imagens de MET (Figura 3e, f). Para o nanocompósito preparado com MWCNTs não modificados (Figura 3e), observou-se uma baixa distribuição nanométrica destas cargas com presença significativa de agregados. No caso de utilização de MWCNTs modificados (Figura 3f), pode ser verificada uma melhor distribuição dos MWCNTs por toda a região do nanocompósito, uma vez que, além de agregados nanométricos, há a presença de nanotubos individuais dispersos ao longo da matriz[16]. Assim, constata-se que o método de preparo dos nanocompósitos a partir de masterbatch propicia uma boa dispersão dos nanotubos de carbono na matriz de PUE e que os MWCNTs modificados apresentam-se como materiais mais indicados para preparação de compósitos com propriedades melhoradas. Segundo a literatura[18], com o tratamento ácido são gerados grupos oxigenados quimissorvidos à superfície dos nanotubos de carbono, como os anidridos de ácido, carbonilas, hidroxilas, grupos carboxílicos, fenólicos e lactônicos, sendo esses três últimos os principais. Os grupos fenólicos, carboxílicos e álcoois são de particular interesse uma vez que podem participar da reação de polimerização com o isocianato originando grupos amida e uretano que 85


Lopes, M. C., Trigueiro, J. P. C., Castro, V. G., Lavall, R. L., & Silva, G. G.

Figura 3. Imagens de MO da superfície dos compósitos contendo 0,5% em massa de: (a) MWCNTs não modificados e (b) MWCNTs modificados. Imagens de MEV da superfície da fratura dos nanocompósitos contendo 0,5% em massa de: (c) MWCNT não modificados e (d) MWCNTs modificados e imagens de MET de cortes de criomicrotomia dos nanocompósitos contendo 0,5% em massa de: (e) MWCNTs não modificados e (f) MWCNTs modificados. 86

Polímeros, 26(1), 81-91, 2016


Otimização do processo de dispersão de nanotubos de carbono em poliuretano termorrígido contribuem para a formação de ligações cruzadas no material final[18]. Diante disso, a espectroscopia vibracional na região do infravermelho foi utilizada para estudar os diferentes materiais. A Figura 4 apresenta o espectro de FTIR do PUE e dos nanocompósitos contendo nanotubos de carbono modificados e não modificados. Através da análise da Figura 4 pode-se observar uma banda larga de baixa intensidade em 3277 cm–1 característica da própria deformação axial da ligação de hidrogênio do grupo N-H com o oxigênio derivado o grupo éter do segmento flexível[16,28,29]. Uma banda na região de 2937 cm–1 corresponde à deformação axial assimétrica de C-H dos grupos CH2[30]. Em 2854 cm–1 foi observada banda de baixa intensidade correspondente à deformação axial de CH2 alifático[30]. Em 1695 cm–1 observa-se banda característica de baixa intensidade correspondente à deformação axial C=O interagindo com ligação de hidrogênio. As bandas agudas que aparecem em 1525 e 1220 cm–1 podem ser atribuídas às ligações C-N de grupos uretanos[30]. A banda em 1101 cm–1 corresponde à deformação axial assimétrica e simétrica da ligação C-O-C[30,31]. Para os nanocompósitos, não foram observadas bandas adicionais ou deslocamentos consideráveis em relação à PUE, nem diferenças devido à funcionalização dos nanotubos. Possivelmente isso está relacionado à baixa concentração de nanotubos utilizada, bem como à presença de apenas 4% de grupos carboxílicos no MWCNT modificado. Dessa forma não foi possível avaliar a diferença entre os MWCNTs, nem se há interação preferencial com os domínios rígidos ou flexíveis do poliuretano termorrígido. A degradação térmica de poliuretanos é um processo complexo e pode ocorrer em pelo menos duas etapas conforme relatado na literatura[32]. A primeira etapa está relacionada à decomposição térmica dos segmentos rígidos e a segunda à degradação dos segmentos flexíveis[32-34]. As curvas TG e DTG para o PUE e os nanocompósitos são apresentadas na Figura 5. Por meio da curva DTG, pode-se observar que a degradação térmica do PUE acontece em pelo menos três estágios, com máximos em 287, 340, 405 °C. A decomposição térmica dos nanocompósitos também ocorreu em três estágios, apresentando pequenas variações de temperatura quando comparados ao PUE. Os resultados mostraram que a estabilidade térmica dos nanocompósitos diminuiu em ~12 e 9 °C para os nanocompósitos contendo MWCNTs não modificados e modificados, respectivamente. O decréscimo na temperatura de degradação para as amostras de MWCNTs não modificados possivelmente ocorreu devido à dispersão dos mesmos na matriz com presença de maiores agregados, observada pelas imagens de microscopias, e à presença de impurezas adsorvidas nas paredes externas dos tubos que contribuíram para um aumento na velocidade de degradação dos segmentos rígidos deste nanocompósito[29,35]. Já para os nanocompósitos contendo MWCNTs modificados, a menor variação provavelmente é devido à melhor dispersão e remoção de impurezas adsorvidas nas superfícies dos tubos como consequência do processo de modificação química. A fim de avaliar a influência dos MWCNTs nas propriedades mecânicas da matriz de PUE, ensaios de tração e dureza foram realizados. Os resultados obtidos para o PUE Polímeros, 26(1), 81-91, 2016

Figura 4. Espectros de FTIR do PUE e dos nanocompósitos contendo 0,5% em massa de MWCNTs modificados e não modificados.

Figura 5. Curva TG e DTG para a PUE e nanocompósitos em atmosfera de N2.

e nanocompósitos produzidos com MWCNTs modificados e não modificados são apresentados na Figura 6. Como já discutido, além das características dos nanotubos de carbono e do procedimento utilizado na dispersão dos mesmos, a natureza da matriz polimérica tem grande influência nas propriedades dos nanocompósitos poliméricos. Portanto, utilizou-se como estratégia para dispersão dos MWCNTs e aumento nas propriedades mecanicas, avaliar a influência do processo pela comparação entre as propriedades dos compósitos preparados com a mesma matriz e nanotubos modificados e não modificados. A tensão na ruptura reduziu cerca de 7% com a introdução de MWCNTs não modificados. Já no nanocompósito produzido a partir de MWCNTs modificados verificou‑se um aumento em torno de 10% em relação à matriz de PUE (Figura 6a). O alongamento na ruptura apresentou um comportamento semelhante, com uma redução de cerca de 3% para o nanocompósito produzido a partir de MWCNTs não modificados e um aumento em torno de 10% para o nanocompósito com MWCNTs modificados, comparando-se aos valores da PUE (Figura 6b). Para o módulo de elasticidade há uma diminuição de cerca de 17% 87


Lopes, M. C., Trigueiro, J. P. C., Castro, V. G., Lavall, R. L., & Silva, G. G. no nanocompósito preparado com MWCNT não modificado e de aproximadamente 8% para o nanocompósito com MWCNT modificado (Figura 6c). Através da Figura 6d, pode-se verificar que os nanotubos de carbono exercem pouca influência na dureza (Shore D) da matriz. Esses resultados corroboram com o que foi observado nos estudos de microscopia, uma vez que mostram que os MWCNTs modificados por tratamento ácido são mais indicados como carga de reforço devido a melhor interação com a matriz e formação de uma interface propícia para a transferência de tensão entre o nanotubo de carbono e o polímero quando o compósito é submetido à tensão mecânica. Esse comportamento foi mostrado na literatura para compósitos preparados com matrizes termorrígidas distintas da utilizada no presente trabalho[18,36]. Para compósitos poliméricos em geral, termoplásticos ou termorrígidos, com MWCNTs, a tendência mais frequente é o aumento no módulo de elasticidade e a diminuição no alongamento de ruptura devido à inserção de uma carga rígida. No entanto, as potenciais influências dos MWCNTs são amplas e também é possível um aumento no alongamento concomitante ao módulo de elasticidade, como já foi relatado para sistemas de PUE/MWCNTs[14], bem como o aumento no alongamento com redução no módulo[37,38]. O efeito nas propriedades é específico para cada sistema,

considerando tipo de nanotubo e tipo de matriz polimérica. Guo e colaboradores relataram uma diminuição do módulo e um aumento do alongamento para compósitos de epóxi (termorrígido de alto módulo) de forma proporcional à concentração utilizada de MWCNTs funcionalizados[38]. Eles discutiram que o aumento na deformação pode ser resultante do deslizamento de camadas internas dos MWCNTs. Portanto, neste caso os nanotubos de carbono atuaram como aditivo para melhorar a ductilidade e não a rigidez do sistema epóxi. Através da funcionalização, esse deslizamento pode ser favorecido devido à camada externa de MWCNTs ser mais defeituosa, a qual interage de forma mais fraca com as camadas internas. Outro fator importante para as propriedades observadas é que a formulação de PUE estudada apresenta um módulo muito elevado, de 290 MPa (valor obtido pela extrapolação da secante a 2% de deformação), enquanto os demais trabalhos na literatura sobre PUE/MWCNTs reportam PUE com módulos abaixo de 10 MPa[14,16,17]. Assim, apesar dos MWCNTs resultarem em redução no módulo, seu valor ainda é mantido em um patamar elevado com a introdução de MWCNTs funcionalizados. A introdução de MWCNTs funcionalizados propiciou um aumento da ductilidade do material e o aumento na tensão na ruptura em 10%, o que é o resultado de ganho

Figura 6. Propriedades mecânicas das amostras estudas. (a) Tensão na ruptura, (b) alongamento percentual na ruptura, (c) módulo de elasticidade (secante a 2%) e (d) dureza (Shore D). 88

Polímeros, 26(1), 81-91, 2016


Otimização do processo de dispersão de nanotubos de carbono em poliuretano termorrígido mais significativo. No caso deste tipo de termorrígido um ganho nesta faixa já é de relevância para considerar-se a aplicação do compósito em projetos de engenharia exigentes. Quando a dispersão de MWCNTs é ineficaz e o compósito é formado por aglomerados grandes, estes atuam como concentradores de tensão, que facilitam a propagação de trincas e reduzem a tensão na ruptura do material[17]. Portanto, somente uma nanocarga especial pode promover algum ganho em propriedades mecânicas para polímeros desta classe. Embora aumentos modestos tenham sido verificados em relação a alguns desses trabalhos, o procedimento descrito permite a preparação de quantidades escaláveis de concentrado. Um diferencial importante para essa rota de processo, que constitui uma maior viabilidade para introdução na indústria, é a utilização de moinho de rolos como equipamento principal de processamento. As referências de PUE/MWCNTs relatam a incorporação da nanocarga por ultrassom ou dispersores de alto cisalhamento[14-16]. No entanto, essas técnicas são baseadas em fontes de dispersão localizadas (nos transdutores ultrassônicos ou nas pontas dos dispersores de alto cisalhamento) e, portanto, apresentam dificuldades no aumento de escala sem resultar em heterogeneidades no processamento. O moinho de rolos garante um processamento homogêneo, onde todo o material é submetido a uma mesma tensão de cisalhamento, sendo uma técnica de alto interesse tecnológico para o aumento de escala. Além disso, os nanotubos de carbono utilizados foram modificados com apenas 4% de grupos carboxílicos (segundo dados do fornecedor e confirmados por TGA). Nosso grupo desenvolveu uma metodologia de funcionalização por tratamento ácido em escala que possibilita a obtenção de nanotubos de carbono com maiores graus de funcionalização sem variação drástica na sua razão de aspecto e com a utilização de quantidades muito pequenas de ácido[21]. Esse procedimento de modificação de superfície dos nanotubos de carbono é muito mais simples dos que as utilizadas nos trabalhos citados anteriormente[37,38]. Assim, acredita-se que a utilização de nanotubos com maior grau de funcionalização aliado ao processo de dispersão otimizado no presente trabalho, trarão efeitos mais significativos nas propriedades mecânicas. A dispersão de nanocargas em polímeros isolantes produz um aumento significativo na condutividade elétrica do material quando a concentração dos nanotubos de carbono atinge o limiar de percolação (aumento de várias ordens de grandeza na condutividade do material em relação à matriz polimérica). Os nanocompósitos obtidos neste trabalho foram estudados por espectroscopia de impedância eletroquímica e a condutividade elétrica foi obtida com o emprego da Equação 1[27,39]. C=

l (1) AxR

sendo l a espessura, A a área e R a resistência dos filmes. A Tabela 1 mostra os resultados de condutividade para todas as amostras investigadas. Os nanocompósitos com MWCNTs modificados e não modificados apresentaram valores de condutividades superiores em aproximadamente duas ordens de grandeza em relação ao valor obtido para a matriz polimérica pura. No entanto, esses valores ainda são menores que Polímeros, 26(1), 81-91, 2016

Tabela 1. Resultados de condutividade para PUE e para os diferentes nanocompósitos preparados. Amostras PUE PUE/MWCNT não modificado 0,5% em massa PU/MWCNT modificado 0,5% em massa

Condutividade (S/cm) 3,4 × 10–6 9,8 × 10–5 8,7 × 10–5

usualmente obtidos considerando-se a alta condutividade dos nanotubos de carbono isolados[40,41]. Segundo Lavall e colaboradores, dependendo da adesão do polímero aos CNTs, há formação de uma camada isolante em torno destas partículas condutoras, diminuindo o contato entre os tubos e, em consequência, as propriedades de condução elétrica[27]. A condutividade de nanocompósitos do tipo polímero/CNTs depende intrinsecamente dessa relação polímero/nanotubo e é sempre menor do que a condutividade obtida pelo contato direto entre os tubos[42]. No presente trabalho, além desse efeito, acredita-se que a concentração utilizada (0,5% de MWCNTs) esteja abaixo do limiar de percolação para esse tipo de sistema. Condutividades entre 10–7 e 10–1 S/cm já foram descritos na literatura para sistemas semelhantes, embora a maioria dos valores estejam na ordem de 10–4 S/cm[42]. Em um trabalho anterior do grupo foram obtidos valores próximos a 10–5 S/cm para a condutividade de um sistema baseado em poliuretano termoplástico (TPU) com 1% em massa de nanotubos modificados[39], mas neste caso o TPU de partida era altamente isolante, apresentando condutividade de 10–12 S/cm. Valores de condutividade da ordem obtida no presente trabalho permitem que o compósito possa ser empregado em aplicações nas quais a dissipação eletrostática seja um requisito[43]. Para atingir condutividades que permitam uso em blindagem eletromagnética deve-se adicionar maior quantidade de nanotubos de carbono, mas será necessário também modificação de estratégia de processamento para produzir dispersões adequadas à percolação elétrica[39,43].

4. Conclusão A produção de masterbatches de MWCNTs em poliol através de uma combinação de misturador de alto cisalhamento e moinho de rolos, possibilitou a fabricação de nanocompósitos com 0,5% em massa de MWCNTs em PUE após diluição, utilizando somente agitação mecânica, sem prejuízo da dispersão dos nanotubos de carbono. Os MWCNTs modificados com grupos oxigenados resultaram em uma melhor dispersão e adesão dos CNTs à matriz polimérica, que refletiram em melhores resultados mecânicos, como tensão na ruptura e alongamento, além de uma melhor preservação da estabilidade térmica. A condutividade dos nanocompósitos foi aumentada em aproximadamente duas ordens de grandeza, o que possibilita o uso do material em aplicações que demandam dissipação eletrostática. A metodologia utilizada garante um processamento homogêneo, com possibilidade de aumento de escala e alta atratividade tecnológica para inserção na indústria. A avaliação do grau de dispersão dos MWCNTs em PUE pelo método desenvolvido neste trabalho e das propriedades obtidas consequentemente, permitem projetar a fabricação de nanocompósitos com diferentes conjuntos de propriedades através do controle do teor de grupos oxigenados nos MWCNTs, da concentração utilizada de nanocarga e das etapas de processo. 89


Lopes, M. C., Trigueiro, J. P. C., Castro, V. G., Lavall, R. L., & Silva, G. G.

5. Agradecimentos Os autores agradecem a Petrobras, Conselho Nacional de Desenvolvimento Científico e Tecnológico (CNPq), Fundação de Amparo à Pesquisa do estado de Minas Gerais (FAPEMIG), Plastiprene Ltda, Instituto Nacional de Ciência e Tecnologia em Nanomateriais de Carbono (INCT) e ao Centro de Microscopia da UFMG. Magnovaldo C. Lopes e João Paulo C. Trigueiro agradecem ao CNPq e a CAPES pelas bolsas concedidas.

6. Referências 1. Prisacariu, C. (2011). Polyurethane elastomers: from morphology to mechanical aspects. New York: Springer. 2. Dieterich, D., Grigat, E., Hahn, W., Hespe, H., & Schmelzer, H. G. (1993). Principles of polyurethane chemistry and special applications. 2. ed. Cincinnati: Hanser. 3. Clemitson, I. R.(2008). Castable polyurethane elastomers. New York: CRC Press. 4. Chattopadhyay, D. K., & Raju, K. V. S. N. (2007). Structural engineering of polyurethane coatings for high performance applications. Progress in Polymer Science, 32(3), 352-418. http://dx.doi.org/10.1016/j.progpolymsci.2006.05.003. 5. Sampurno, Y., Borucki, L., Zhuang, Y., Misra, S., Holland, K., Boning, D., & Philipossian, A. (2009). Characterization of thermoset and thermoplastic polyurethane pads, and molded and non-optimized machined grooving methods for oxide chemical mechanical planarization applications. Thin Solid Films, 517(5), 1719-1726. http://dx.doi.org/10.1016/j. tsf.2008.09.077. 6. Ledru, Y., Bernhart, G., Piquet, R., Schmidt, F., & Michel, L. (2010). Coupled visco-mechanical and diffusion void growth modelling during composite curing. Composites Science and Technology, 70(15), 2139-2145. http://dx.doi.org/10.1016/j. compscitech.2010.08.013. 7. Wang, T.-L., & Tseng, C.-G. (2007). Polymeric carbon nanocomposites from multiwalled carbon nanotubes functionalized with segmented polyurethane. Journal of Applied Polymer Science, 105(3), 1642-1650. http://dx.doi. org/10.1002/app.26224. 8. Wei, B. Q., Vajtai, R., & Ajayan, P. M. (2001). Reliability and current carrying capacity of carbon nanotubes. Applied Physics Letters, 79(8), 1172-1174. http://dx.doi.org/10.1063/1.1396632. 9. Yang, D. J., Wang, S. G., Zhang, Q., Sellin, P. J., & Chen, G. (2004). Thermal and electrical transport in multi-walled carbon nanotubes. Physics Letters, Part A, 329(3), 207-213. http://dx.doi.org/10.1016/j.physleta.2004.05.070. 10. Yu, M.-F., Lourie, O., Dyer, M. J., Moloni, K., Kelly, T. F., & Ruoff, R. S. (2000). Strength and breaking mechanism of multiwalled carbon nanotubes under tensile load. Science, 287(5453), 637-640. http://dx.doi.org/10.1126/science.287.5453.637. PMid:10649994. 11. Moniruzzaman, M., & Winey, K. I. (2006). Polymer nanocomposites containing carbon nanotubes. Macromolecules, 39(16), 5194-5205. http://dx.doi.org/10.1021/ma060733p. 12. Silva, M. A., Tavares, M. I. B., Nascimento, S. A. M., Rodrigues, E., & Jd, R. (2012). Caracterização de nanocompósitos de poliuretano/montmorilonita organofílica por RMN de baixo campo. Polímeros: Ciência e Tecnologia, 22(5), 481-485. http://dx.doi.org/10.1590/S0104-14282012005000064. 13. De Volder, M. F. L., Tawfick, S. H., Baughman, R. H., & Hart, A. J. (2013). Carbon Nanotubes: Present and Future Commercial Applications. Science, 339(6119), 535-539. http:// dx.doi.org/10.1126/science.1222453. PMid:23372006. 90

14. Kantheti, S., Gaddam, R. R., Narayan, R., & Raju, K. V. (2014). Hyperbranched polyol decorated carbon nanotube by click chemistry for functional polyurethane urea hybrid composites. RSC Advances, 4(47), 24420-24427. http://dx.doi. org/10.1039/c4ra02442g. 15. McClory, C., McNally, T., Brennan, G. P., & Erskine, J. (2007). Thermosetting polyurethane multiwalled carbon nanotube composites. Journal of Applied Polymer Science, 105(3), 1003-1011. http://dx.doi.org/10.1002/app.26144. 16. Xiong, J., Zheng, Z., Qin, X., Li, M., Li, H., & Wang, X. (2006). The thermal and mechanical properties of a polyurethane/ multi-walled carbon nanotube composite. Carbon, 44(13), 2701-2707. http://dx.doi.org/10.1016/j.carbon.2006.04.005. 17. Xia, H., & Song, M. (2006). Preparation and characterisation of polyurethane grafted single-walled carbon nanotubes and derived polyurethane nanocomposites. Journal of Materials Chemistry, 16(19), 1843-1851. http://dx.doi.org/10.1039/ b601152g. 18. Karabanova, L. V., Whitby, R. L. D., Korobeinyk, A., Bondaruk, O., Salvage, J. P., Lloyd, A. W., & Mikhalovsky, S. V. (2012). Microstructure changes of polyurethane by inclusion of chemically modified carbon nanotubes at low filler contents. Composites Science and Technology, 72(8), 865-872. http:// dx.doi.org/10.1016/j.compscitech.2012.02.008. 19. Lopes, M. C., Silva, G. G., Lavall, R. L., Diniz, V. P. A., & Castro, V. G. (2013). BR Patente No 10201330082961. Brasília: Instituto Nacional de Propriedade Industrial. 20. Lopes, M. C., Castro, V.G., Seara, L. M., Diniz, V. P. A., Lavall, R. L., & Silva, G. G. (2014). Thermosetting polyurethanemultiwalled carbon nanotube composites: thermomechanical properties and nanoindentation. Journal of Applied Polymer Science, 131(23) 41207 21. Silva, G. G., Lavall, R. L., Figueiredo, K. C. S., Castro, V. G., Costa, I. B., Medeiros, F. S., Lopes, M. C., Ferreira, F. L. Q., & Diniz, V. P. A. (2014). BR Patente No 1020140259660. Brasília: Instituto Nacional de Propriedade Industrial. 22. Delpech, M. C., Coutinho, F. M. B., Sousa, K. G. M., & Cruz, R. C. (2007). Estudo Viscosimétrico de Prepolímeros Uretânicos. Polímeros: Ciência e Tecnologia, 17(4), 294-298. http://dx.doi.org/10.1590/S0104-14282007000400008. 23. Pacheco, M. F. M., Fiorio, R., Zattera, A. J., Zeni, M., & Crespo, J. S. (2007). Efeito da concentração de segmentos rígidos nas propriedades físico-mecânicas, químicas e na morfologia de elastômeros microcelulares de poliuretano. Polímeros: Ciência e Tecnologia, 17(3), 234-239. http://dx.doi. org/10.1590/S0104-14282007000300013. 24. Ryszkowska, J. (2009). Quantitative image analysis of polyurethane/carbon nanotube composite microstructures. Materials Characterization, 60(10), 1127-1132. http://dx.doi. org/10.1016/j.matchar.2009.01.021. 25. Ryszkowska, J., Jurczyk-Kowalska, M., Szymborski, T., & Kurzydlowski, K. J. (2007). Dispersion of carbon nanotubes in polyurethane matrix. Physica E: Low-Dimensional Systems and Nanostructures, 39(1), 124-127. http://dx.doi.org/10.1016/j. physe.2007.02.003. 26. Ma, P.-C., Siddiqui, N. A., Marom, G., & Kim, J.-K. (2010). Dispersion and functionalization of carbon nanotubes for polymer-based nanocomposites: a review. Composites Part A: Applied Science and Manufacturing, 41(10), 1345-1367. http://dx.doi.org/10.1016/j.compositesa.2010.07.003. 27. Lavall, R. L., Sales, J. A., Borges, R. S., Calado, H. D. R., Machado, J. C., Windmoller, D., Silva, G. G., Lacerda, R. G., & Ladeira, L. O. (2010). Nanocompósitos de poliuretana termoplástica e nanotubos de carbono de paredes múltiplas para dissipação eletrostática. Quimica Nova, 33(1), 133-140. http://dx.doi.org/10.1590/S0100-40422010000100025. Polímeros, 26(1), 81-91, 2016


Otimização do processo de dispersão de nanotubos de carbono em poliuretano termorrígido 28. Wen, T.-C., Du, Y.-L., & Digar, M. (2002). Compositional effect on the morphology and ionic conductivity of thermoplastic polyurethane based electrolytes. European Polymer Journal, 38(5), 1039-1048. http://dx.doi.org/10.1016/S0014-3057(01)00257-9. 29. Xiong, J., Zheng, Z., Song, W., Zhou, D., & Wang, X. (2008). Microstructure and properties of polyurethane nanocomposites reinforced with methylene-bis-ortho-chloroanilline-grafted multi-walled carbon nanotubes. Composites. Part A, Applied Science and Manufacturing, 39(5), 904-910. http://dx.doi. org/10.1016/j.compositesa.2007.12.008. 30. Wang, T.-L., Yu, C.-C., Yang, C.-H., Shieh, Y.-T., Tsai, Y.-Z., & Wang, N.-F. (2011). Preparation, characterization, and properties of polyurethane-grafted multiwalled carbon nanotubes and derived polyurethane nanocomposites. Journal of Nanomaterials, 2011, 1-9. http://dx.doi.org/10.1155/2011/814903. PMid:21808638. 31. Barick, A. K., & Tripathy, D. K. (2011). Preparation, characterization and properties of acid functionalized multiwalled carbon nanotube reinforced thermoplastic polyurethane nanocomposites. Materials Science and Engineering B, 176(18), 1435-1447. http://dx.doi.org/10.1016/j.mseb.2011.08.001. 32. Chattopadhyay, D. K., & Webster, D. C. (2009). Thermal stability and flame retardancy of polyurethanes. Progress in Polymer Science, 34(10), 1068-1133. http://dx.doi.org/10.1016/j. progpolymsci.2009.06.002. 33. Hablot, E., Zheng, D., Bouquey, M., & Avérous, L. (2008). Polyurethanes based on castor oil: kinetics, chemical, mechanical and thermal properties. Macromolecular Materials and Engineering, 293(11), 922-929. http://dx.doi.org/10.1002/ mame.200800185. 34. Javni, I., Petrović, Z. S., Guo, A., & Fuller, R. (2000). Thermal stability of polyurethanes based on vegetable oils. Journal of Applied Polymer Science, 77(8), 1723-1734. http:// dx.doi.org/10.1002/1097-4628(20000822)77:8<1723::AIDAPP9>3.0.CO;2-K. 35. Mondal, S., & Hu, J. L. (2006). Thermal degradation study of functionalized MWNT reinforced segmented polyurethane membrane. Journal of Elastomers and Plastics, 38(3), 261-271. http://dx.doi.org/10.1177/0095244306064237. 36. Karabanova, L., Whitby, R. D., Bershtein, V., Korobeinyk, A., Yakushev, P., Bondaruk, O., Lloyd, A., & Mikhalovsky,

Polímeros, 26(1), 81-91, 2016

S. (2013). The role of interfacial chemistry and interactions in the dynamics of thermosetting polyurethane-multiwalled carbon nanotube composites at low filler contents. Colloid & Polymer Science, 291(3), 573-583. http://dx.doi.org/10.1007/ s00396-012-2745-4. 37. Zhao, C., Ji, L., Liu, H., Hu, G., Zhang, S., Yang, M., & Yang, Z. (2004). Functionalized carbon nanotubes containing isocyanate groups. Journal of Solid State Chemistry, 177(12), 4394-4398. http://dx.doi.org/10.1016/j.jssc.2004.09.036. 38. Song, H.-J., Zhang, Z.-Z., & Men, X.-H. (2007). Surfacemodified carbon nanotubes and the effect of their addition on the tribological behavior of a polyurethane coating. European Polymer Journal, 43(10), 4092-4102. http://dx.doi.org/10.1016/j. eurpolymj.2007.07.003. 39. Lima, A. M. F., Castro, V. G. D., Borges, R. S., & Silva, G. G. (2012). Electrical conductivity and thermal properties of functionalized carbon nanotubes/polyurethane composites. Polímeros Ciência e Tecnologia, 22(2), 117-124. http://dx.doi. org/10.1590/S0104-14282012005000017. 40. Ajayan, P. M. (1999). Nanotubes from carbon. Chemical Reviews, 99(7), 1787-1800. http://dx.doi.org/10.1021/cr970102g. PMid:11849010. 41. Barrau, S., Demont, P., Peigney, A., Laurent, C., & Lacabanne, C. (2003). DC and AC conductivity of carbon nanotubespolyepoxy composites. Macromolecules, 36(14), 5187-5194. http://dx.doi.org/10.1021/ma021263b. 42. Musumeci, A. W., Silva, G. G., Liu, J.-W., Martens, W. N., & Waclawik, E. R. (2007). Structure and conductivity of multiwalled carbon nanotube/poly(3-hexylthiophene) composite films. Polymer, 48(6), 1667-1678. http://dx.doi.org/10.1016/j. polymer.2007.01.027. 43. Ramasubramaniam, R., Chen, J., & Liu, H. (2003). Homogeneous carbon nanotube/polymer composites for electrical applications. Applied Physics Letters, 83(14), 2928-2930. http://dx.doi. org/10.1063/1.1616976. Enviado: Maio 31, 2015 Revisado: Ago. 13, 2015 Aceito: Set. 02, 2015

91


http://dx.doi.org/10.1590/0104-1428.1920

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

Investigação do efeito do tempo de exposição à temperatura ambiente e ao tempo de estocagem de um filme adesivo estrutural de resina epoxídica Investigation of the effect of exposure time at room temperature in a structural epoxy resin adhesive film Ana Carolina Teixeira Neves da Silva1, Fernanda Guilherme1, Vanesa Mitchell Ferrari1* e Paulo Eduardo Ferrari2 1

Faculdade de Engenharias, Arquitetura e Urbanismo – FEAU, Universidade do Vale do Paraíba – UNIVAP, São José dos Campos, SP, Brasil 2 Embraer S.A., São José dos Campos, SP, Brasil *vanesa_mitchell@yahoo.com.br

Resumo A exposição de filmes adesivos de resina epoxídica à temperatura ambiente propicia condições que afetam diretamente suas propriedades físicas. Para o adesivo estrutural AF191, de aplicação no setor aeronáutico, isto pode ser confirmado pela súbita perda de pegajosidade (“tack”). Sendo assim, o comportamento do sistema de resina escolhido foi estudado experimentalmente, expondo-se este ao ambiente de laboratório a intervalos pré-determinados de tempo e a diferentes períodos de armazenamento (–18°C). O efeito da exposição foi analisado por meio de ensaios fisico-químicos baseados em normas internacionais complementados por testes empíricos não normatizados que visaram observar alguns comportamentos especificos da resina, como capacidade de fluxo e solubilidade. A partir dos resultados, foi possível observar alterações no comportamento fisico-químico, que afetam as condições de processabilidade do adesivo, como a capacidade de fluxo do adesivo quando aquecido, redução da aderência e da solubilidade em solvente orgânico. Estas alterações sugerem que houve um possível aumento no número de reticulações da matriz polimérica à temperatura ambiente.Porém este fato não foi comprovadopelo método de análise térmica utilizado (calorimetria exploratória diferencial – DSC), onde não foi observado alteração significativa da entalpia de cura ao longo do período de exposição. Isso pode ser atribuído a uma taxa de cura muito baixa da matriz ou apenas pela falta sensibilidade do método. A mudança de comportamento pode estar associada ao processo de plastificação da matriz polímérica causado pela absorção de umidade devido ao caráter higroscópico da matriz epoxídica, que pode ser verificado no aumento progressivo do ensaio de teor de voláteis. O período de estocagem mostrou que não houve alteração significativa das características avaliadas do adesivo. Palavras-chaves: adesivo, resina epoxídica, cura, fluidez, solubilidade. Abstract The exposure of epoxy based adhesive films to room temperature provides conditions that directly affects their physical properties. For structural adhesive AF191, application in the aerospace industry, this can be confirmed by the sudden loss of tack. Therefore, the behavior of the chosen resin system was experimentally studied, by means of the exposure to laboratory environment at pre determinated time interval and different storage periods (–18°C).The effect of the exposure was evaluated by means of physicochemical tests based on international standards, being complimented by not normalized empirical tests that aimed to observe some specific behavior of the resin, like flow and solubility. From the results, it was possible to observe changes in the physicochemical behavior that affects the adhesive processability conditions such as adhesive flow capacity, self adhesiveness reduction and solubility to organic solvent. These changes suggest that there was a possible increase in the number of crosslinking of the polymer matrix at room temperature. This fact however was not proven by the used thermal analysis method (differential scanning calorimetry – DSC), where since there was no significant change in cure enthalpy was observed during the exposure period. This can be attributed to a very low curing rate of the matrix or just low sensitivity of the method. The change in behavior can be related to the lamination process the polymer matrix caused by moisture absorption due to the hygroscopic nature of the epoxy matrix, which can be found in the progressive increase in the volatile content test. The storage period was efficient indicating no significant change in the characteristics evaluated the adhesive. Keywords: adhesive, epoxy resin, cure, flow, solubility.

92

Polímeros, 26(1), 92-100, 2016


Investigação do efeito do tempo de exposição à temperatura ambiente e ao tempo de estocagem de um filme adesivo estrutural de resina epoxídica

1. Introdução O termo adesivo é utilizado para descrever a substância capaz de manter materiais unidos pelo contato entre superfícies. O adesivo estrutural é definido como agente de ligação requerido para transferir cargas necessárias entre superfícies (aderentes) expostas ao ambiente de trabalho típico para a estrutura envolvida[1]. O termo adesivo estrutural costuma ser utilizado para definir um adesivo de resistência elevada de forma a se atingir o sucesso da montagem em seu ambiente operacional com alto grau de confiabilidade, como os filmes epoxídicos e acrílicos. Já os adesivos não estruturais, como a cola de madeira e selantes, possuem menor resistência, e/ou geralmente são usados para fixação temporária[2]. Atualmente, existem diversas técnicas utilizadas para unir duas superfícies como, por exemplo, a soldagem, prendedores mecânicos (rebites, parafusos, etc) e colagem. A utilização dos adesivos minimiza os efeitos de concentração de tensão, aumenta a flexibilidade da peça e age como vedação a intempéries. Na área de manufatura, a utilização dos adesivos estruturais para colagem elimina a necessidade de furos, reduz o peso e o custo total do processo (Figura 1). Normalmente, a escolha pelo processo e o uso de adesivos pode envolver os estudos preliminares de desempenho, da capacidade de produção, de custo e confiabilidade[4]. Os filmes adesivos são, na atualidade, os materiais que despertam maior interesse nos estudos de propriedades e na evolução da técnica denominada colagem[3]. Os filmes adesivos estruturais de resina epoxídica, largamente utilizados para aplicações na indústria aeronáutica, são geralmente comercializados como filmes finos e contínuos, que devem ser armazenados sob refrigeração (geralmente -18oC). Estes são preferidos aos adesivos líquidos e pastas devido à sua uniformidade e reduzido conteúdo de vazios. Os filmes adesivos possuem agentes que requerem alta temperatura de cura e os mantêm relativamente estáveis à temperatura ambiente por períodos curtos de tempo[5]. As resinas epoxídicas são preferidas porque propiciam a produção de filmes adesivos que, depois de curados, apresentam ótima resistência a solventes, combustíveis e lubrificantes. São amplamente utilizados na indústria aeronáutica, tanto na colagem de materiais metálicos quanto compósitos, devido à diversidade de suas propriedades, garantindo elevada relação entre custo, peso e resistência química a diversos ambientes operacionais, além de excelentes propriedades térmicas, físicas e mecânicas[1,5].

Figura 1. Esquematização de uma junção de peças através de (A) rebites, (B) adesivos estruturais[3]. Polímeros, 26(1), 92-100, 2016

Os filmes adesivos epoxídicos apresentam aderência à maioria dos materiais sólidos, independentemente da forma, espessura ou de suas propriedades físicas. Apesar de apresentarem custo elevado em nível de matériaprima, proporcionam grande redução de peso e eliminam operações posteriores de acabamento, além da facilidade na sua montagem. Os adesivos podem ser produzidos em várias formas: soluções à base de solventes, emulsão aquosa, filme suportado ou filme não suportado[4]. No processo de fabricação do filme adesivo em estudo, a resina em sua fase líquida impregna totalmente fibras de poliésteres, poliamidas (nylon) ou fibras de vidro. Adesivos com este tipo de construção facilitam o manuseio antes da cura e promovem melhor controle de fluxo de adesivo durante esta, auxiliando o controle de espessura da área de colagem[5]. O processo de fabricação dos filmes adesivos é feito de forma contínua e altamente automatizada, onde as fibras de suporte, em forma de uma manta fina, são depositadas sobre um filme de resina (matriz) e passada entre mandris que regulam a distribuição e homogeneidade da resina enquanto fluída. Esse sistema é então passado em regiões de aquecimento para que a resina atinja um estagio de polimerização intermediária, onde esta não escorre mais, porém, não está completamente curada, denominado estágio B. Após essa etapa, o material é coberto por filme plástico protetivo de polietileno e silicone que servem de separador e como proteção contra contaminação e umidade. Finalmente, os filmes são enrolados em um suporte e armazenados em forma de bobinas[6,7]. No momento da utilização, o filme adesivo é retirado do armazenamento, descongelado até a temperatura ambiente ainda na embalagem, desembalado e cortado no formato desejado, sendo o plástico de proteção retirado no momento da aplicação[3]. A forma de filme garante um peso ideal e proporções controladas e exatas de resina e catalisador no adesivo. Durante o ciclo de aquecimento, o filme se liquefaz e flui o suficiente para molhar as superfícies a serem coladas, permitindo a remoção do ar retido e passando pelo processo de cura, quando se transformam em um sólido infusível[8]. No processo de cura ocorre a formação de ligações cruzadas covalentes entre as cadeias moleculares adjacentes. Estas ligações ancoram as cadeias entre si para resistir aos movimentos vibracionais e rotacionais a altas temperaturas. Apenas o aquecimento até temperaturas excessivas causará a quebra destas ligações onde ocorrerá a degradação do polímero[8]. O tempo de armazenamento (“shelflife”) para este tipo de material, que é considerado perecível, é limitado pelo prazo de validade do material em condições de estocagem recomendadas (geralmente –18°C), com o intuito de preservar ao máximo as propriedades químicas e físicas adequadas ao seu uso, contabilizado pelo período de tempo a partir da sua produção. Este tempo pode variar de poucos dias a mais de um ano, dependendo da natureza do material ou do seu manuseio[2]. Outro controle utilizado para este tipo de material é o tempo acumulado do material quando retirado do ambiente de armazenamento e exposto ao ambiente de produção, (“out time”). 93


Silva, A. C. T. N., Guilherme, F., Ferrari, V. M., & Ferrari, P. E. Nos últimos anos, vários artigos relataram o envelhecimento físico em resinas epoxídicas[9]. A influência do envelhecimento físico nas propriedades viscoelásticas das resinas baseadas no DGEBA foi bastante estudada por McKenna e colaboradores, pois esta resina serve como modelo no estudo das estruturas reticuladas[10]. As alterações estruturais produzidas devido ao envelhecimento físico dependem do tempo e da temperatura e evidenciam o comportamento dos materiais termofixos a temperaturas abaixo da transição vítrea sendo, normalmente, o resultado do relaxamento lento do material vítreo a partir do seu estado inicial de não equilíbrio termodinâmico para o seu estado final de equilíbrio[9]. Todas as propriedades que dependem do volume específico (ou do volume livre) se alteram durante este processo e, portanto, o envelhecimento físico possui importância prática na determinação das propriedades de engenharia destes materiais[10]. No trabalho de Kim e co-autoresfoi investigada a influência das condições ambientais às quais o prepreg de resina epoxídica é submetido antes do processo de cura para materiais curados fora da autoclave.Verificou-se por meio de análise MDSC e reometria que a exposição do material à temperaturaambiente alterou o grau inicial de cura da resina, bem como a evolução da taxa de cura e a viscosidade a temperaturas elevadas. Por meio do acompanhamento da evolução das propriedades dielétricas da resina por análise DEA, em condições ambientee a temperatura constante, foram desenvolvidas correlações eficazes para se obter o grau de cura e viscosidade. Comprova-se a eficácia do método de acompanhamento das propriedades dielétricas para o estudo do processo de cura, identificando propriedades e transições da resina e, além disso, prevendo a influência das condições ambientais no processo de polimerização[11]. O envelhecimento dos adesivos pode ser relacionado com a temperatura e umidade, e mesmo à exposição à luz. Estes efeitos são sempre prejudiciais e a taxa de deterioração depende do tipo e forma da resina[2]. A umidade adquirida pela resina epoxídica pode se apresentar de diferentes formas, parte das moléculas de água absorvidas podem se difundir entre as cadeias do polímero, atuar como plastificante aumentando o volume livre e consequentemente, reduzir a temperatura de transição vítrea, outra parte das moléculas de água podem se alojar no interior de microcavidades da matriz polimérica ou se difundir na região da interfase fibra-matriz[12-14], e ainda existe a possibilidade de participar de reações de hidrólise[13]. Assim, dependendo da forma que estas moléculas de água se apresentam é possível que causem diferentes efeitos sobre o material[13]. Sales, por meio das técnicas de luminescência e TG, mostrou que poderia ter ocorrido uma plastificação da matriz polimérica devido à presença de água. Embora tenha apresentado bons resultados, a técnica de luminescência é disponível somente em poucos laboratórios especializados, sendo assim, nesse atual trabalho foram usadas técnicas mais convencionais[15].

A exposição de adesivos ao meio ambiente ao longo do tempo provoca alterações de características como aumento da viscosidade e a redução de pegajosidade (”tack”) que durante o processo de laminação compromete a capacidade do adesivo de se conformar a raios e contornos (”drape”) durante o processo de moldagem. O tempo de gelificação (”gel time”) diminui e ocorre perda da capacidade de fluxo durante a cura, prejudicando potencialmente o tempo para a remoção de voláteis durante o processo de cura[16]. Assim, o presente estudo, faz uso de metodologias para investigar o comportamento da resina tanto quando exposta à temperatura ambiente quanto armazenada em freezer por longos períodos, assim como a combinação dessas condições.

2. Experimental O material em estudo é um filme adesivo estrutural de uso aeronáutico, tendo como componente predominante na sua formulação uma matriz de resina epoxídica sustentada por uma trama de fibra de poliamida, denominado comercialmente como AF191, produzido pela empresa 3M. Para as análises, foram utilizados os seguintes equipamentos: - Calorímetro Exploratório Diferencial, DSC Perkin Elmer – Modelo Pyris1; - Equipamento de Gel Time – Modelo Fishers Johns Melting Point.

3. Caracterização do Material Para estudaro efeito do tempo de exposição iniciou-se a preparação dos corpos de provas (cdp´s) com o corte do filme em ambiente controlado de temperatura (22°C ± 3°C) e umidade relativa do ar (50% ± 5%) e submetidos a essa condição ambiente em diferentes intervalos de tempo, contabilizados em “dias de exposição”. A condição do material no momento do corte das amostras foi considerada como “zero dia de exposição”. Para o processo de envelhecimento à temperatura ambiente, as amostras foram submetidas a diferentes intervalos de tempo (3, 5, 8, 10 e 12 dias). Ao atingirem o tempo de exposição desejado, as amostras foram embaladas em filme plástico, seladas, identificadas e armazenadas em freezer à –18 °C (Tabela 1). Para verificar o efeito do tempo de exposição e do tempo estocagem do material no freezer simultaneamente, foram preparados 4 conjuntos de amostras (kits), cada um possuindo amostras submetidas a diferentes intervalos de tempo de exposição à temperatura ambiente (3, 5, 8, 10 e 12 dias). Cada kit foi armazenado em freezer por quatro intervalos de tempo diferentes (0, 4, 8 e 12 meses), conforme indicado na Tabela 1.

Tabela 1. Descrição das amostras ensaiadas por Kit, dias de exposição e tempo de estocagem. Exposição ambiente (dias) Estocagem em freezer (meses)

94

Kit 1 3, 5, 8, 10 e 12 0

Kit 2 3, 5, 8, 10 e 12 4

Kit3 3, 5, 8, 10 e 12 8

Kit 4 3, 5, 8, 10 e 12 12

Polímeros, 26(1), 92-100, 2016


Investigação do efeito do tempo de exposição à temperatura ambiente e ao tempo de estocagem de um filme adesivo estrutural de resina epoxídica Os cdp´s preparados foram submetidos aos ensaios de tempo de gelificação, teor de voláteis, capacidade de fluxo, aderência e solubilidade. O ensaio de tempo de gelificação é baseado na norma ASTM D3532-M[17]. Consiste em determinar o tempo em que, a uma determinada temperatura, a resina polimérica passa do estado líquido para o início do estado sólido, determinando a perda da fluidez da matriz. Para tal, o equipamento Gel Time é aquecido na faixa de temperatura de 166,4°C a 177,6°C. Coloca-se então uma lâmina de vidro sobre o equipamento para aquecimento por 20 segundos. Em seguida adiciona-se a amostra e uma segunda lâmina de vidro. O tempo é cronometrado até o ponto de gelificação da resina, considerado como o momento de transformação para o estado sólido. Para a análise de teor de voláteis, as amostras foram cortadas nas dimensões de 10 cm2, pesadas em balança analítica e colocadas na estufa na faixa de temperatura de 166,4 °C a 177,6°C por 1 hora. Após este período, foram retiradas da estufa e colocadas em dessecador até atingirem a temperatura ambiente. Uma vez estabilizada a temperatura, são novamente pesadas e calcula-se a perda de massa conforme ASTM D3530/D3530M–97[18]. O ensaio de capacidade de fluxo foi realizado para determinar o teor de fluxo sob pressão e temperatura controlada[19]. Para o ensaio foram cortadas amostras com o diâmetro de 6,5 cm. Um tecido de teflon não poroso foi colocado sobre o adesivo e o conjunto posicionado em prensa aquecido na faixa de temperatura de 166,4°C a 177,6°C em prensa, durante 1 hora. O valor de fluxo é expresso em termos de variação percentual de área. Para o presente trabalho, a variação da área foi calculada a partir do diâmetro inicial da amostra e de um diâmetro final médio aproximado (D1), calculado por meio do comprimento (L) e largura máxima(W) resultantes após aquecimento (Figura 2). A capacidade de aderência é avaliada por meio do contato do material não curado a si mesmo e a uma superfície

metálica posicionada na vertical, em condição ambiente. O ensaio adotado foi baseado na norma NMS 5320[20], o qual inicia-se com o corte das amostras nas dimensões de 10 cm × 5 cm. A primeira amostra é fixada em uma chapa de alumínio de 400 cm2 com uma leve pressão exercida por meio de uma espátula plástica. Posteriormente, aplica-se sobre a amostra uma segunda camada para verificar a aderência do filme a si mesmo. A placa é posicionada verticalmente durante 30 minutos. Após esse período, é definido o nível de aderência com base no critério definido na Tabela 2. A aderência é um fator importante para estes adesivos, pois determina a qualidade da colagem em um substrato ou sobre o próprio adesivo[20]. Não foi encontrado na literatura um método para determinar a solubilidade de adesivos em solvente orgânico, portanto uma metodologia específica foi proposta para avaliar a solubilidade através da extração da resina pelo solvente. Neste caso, foi utilizado um solvente orgânico à base de éster dibásico, conhecido comercialmente como Rhodiasolv (Rhodia), com alto poder de solvência. É um produto não tóxico e não inflamável, biodegradável e de baixo VOC (composto orgânico volátil)[21]. Para a realização do ensaio, as amostras foram cortadas com 50 cm2, em formato retangular (10 cm × 5cm), pesadas em balança analítica e imersas uma a uma em um béquer com 300 ml do solvente, com agitação constante, durante 30 segundos. Em seguida, a amostra é retirada e deixada em descanso até a evaporação total do resíduo de solvente, sendo pesada novamente para a realização do cálculo da perda de material. Os ensaios em DSC foram realizados conforme método dinâmico nas amostras do kit 3, com variação de temperatura de 25°C a 350°C, com uma razão de aquecimento de 10°C/min e atmosfera inerte de N2 com vazão 20 ml/min. Foram utilizados aproximadamente 8,5 mg do material por amostra. Este método de caracterização foi utilizado como técnica auxiliar para avaliar termicamente a cura do filme adesivo. Os parâmetros investigados foram a temperatura onset (Tonset), a temperatura no máximo do pico (Tmáx) e a variação de entalpia (ΔE)[22].

4. Resultados e Discussão A seguir estão apresentados os resultados dos ensaios físico-químicos adotados.

4.1 Gelificação Figura 2. Modelo de medição de área do fluxo, após a cura do adesivo. (L: medida do comprimento máximo, W: medida da largura, essas medidas definem D1) e fórmula para cálculo.

Os valores encontrados por meio deste ensaio (Tabela 3) indicam que quanto maior o tempo de exposição, menor o tempo de gelificação do polímero, o que pode ser

Tabela 2. Níveis de Aderência[20]. Nível I Nível II Nível III Nível IV Nível V Nível VI

Baixa aderência. Material seco e rígido. Baixa aderência, porém maleável Leve aderência a si mesmo. Incapaz de aderir à superfície da placa metálica na vertical por mais de 30 minutos; Boa aderência a si mesmo. Capaz de aderir à superfície da placa metálica na vertical por mais de 30 minutos; Muita aderência; A resina cola nas luvas. Capaz de aderir a si mesmo e à superfície da placa metálica na vertical por mais de 30 minutos; Alta aderência. Material “grudento e molhado”. A resina cola nas luvas. Existe transferência de resina. Capaz de aderir a si mesmo e à superfície da placa metálica na vertical por mais de 30 minutos;

Polímeros, 26(1), 92-100, 2016

95


Silva, A. C. T. N., Guilherme, F., Ferrari, V. M., & Ferrari, P. E. verificado pela queda gradual e contínua do tempo de gelificação ao longo dos 12 dias de exposição para todos os kits, provavelmente devido a um ligeiro aumento do número de ligações cruzadas do polímero termorrígido ocorrido mesmo à temperatura ambiente. Comparando-se os tempos de gelificação entre os diferentes kit para verificar a eficiência de estocagem do material em freezer, tomando o kit 1(tempo de estocagem igual a 0 meses), observa-se uma redução no tempo de gelificação quando o material é novo (zero dias de exposição)e a partir de 10 dias de exposição à temperatura ambiente (10 e 12 dias de exposição) quando comparado aos demais kits. Este processo de reticulação foi evitado durante o período de estocagem em freezer porque não houve alteração significativa nos tempo de gelificação entre os kits 2, 3 e 4.

4.2 Capacidade de fluxo A Tabela 4 apresenta os resultados dos ensaios de capacidade de fluxo para os quatro ciclos de armazenamento. Apesar do elevado desvio padrão apresentado nos resultados encontrados neste ensaio é possível observar

uma diminuição progressiva na capacidade do adesivo em fluir com o aumento do tempo de exposição ao ambiente em todos os kits, sugerindo uma possível reticulação entre as cadeias poliméricas. Assim, essa variação pode ser considerada como inerente à metodologia de ensaio. Comparando-se os resultados entre os kits não é possível afirmar que o período de estocagem tenha influenciado na capacidade de fluxo do material. Na Figura 3 estão apresentados os cdp’s resultantes do ensaio para as amostras do Kit 1. Pode-se observar claramente a redução da capacidade de fluxo pela redução da área após o escoamento da resina por meio da modificação das variáveis, L(medida do comprimento máximo), W (medida da largura), e consequentemente a variação de D1.

4.3 Aderência Os resultados dos ensaios de aderência para os quatro ciclos de armazenamento estão apresentados na Tabela 5: Pode-se observar que para as amostras do kit 1, nos períodos de exposição iniciais (0 e 3 dias) não houve modificação considerável de aderência sendo notado variação

Tabela 3. Tempos de gelificação (min). Exposição (dias) 0 3 5 8 10 12

Kit 1 13,82 ± 0,55 9,23 ± 0,25 8,53 ± 0,35 8,22 ± 0,15 7,58 ± 0,22 5,93 ± 0,25

Kit 2 11,12 ± 0,18 9,53± 0,23 8,6 ± 0,15 8,42 ± 0,12 5,83± 0,22 2,98 ± 0,17

Kit3 10,55 ± 0,57 10,63 ± 33 9,1 ± 0,30 8,87 ± 0,20 5,27 ± 0,55 3,05 ± 0,60

Kit 4 10, 08 ± 0,3 9,72 ± 0,37 8,88 ± 0,32 7,9 ± 0,27 6,00± 33 s 2,98 ± 0,32

Kit 2 368 ± 14 330 ± 14 268 ± 20 153 ± 8 117 ± 6 77 ± 11

Kit 3 391 ± 7 343 ± 25 270 ± 36 201 ± 4 154 ± 40 78 ± 7

Kit 4 367 ± 32 345 ± 25 311 ± 6 273 ± 31 215 ± 72 93 ± 10

Tabela 4. Capacidade de fluxo (%). Exposição (dias) 0 3 5 8 10 12

Kit 1 399 ± 14 345 ± 15 299 ± 16 155 ± 4 157 ± 4 76 ± 6

Figura 3. Fotos do ensaio de fluxo, o número indicado em cada foto corresponde ao período de envelhecimento de cada material em dias. 96

Polímeros, 26(1), 92-100, 2016


Investigação do efeito do tempo de exposição à temperatura ambiente e ao tempo de estocagem de um filme adesivo estrutural de resina epoxídica detectável de aderência (de nível IV para III) apenas após o quinto dia de exposição, permanecendo nesse nível até 12 dias. Quando comparados os dados referentes aos tempos de exposição e de armazenamento, nota-se certa influência também do tempo de armazenamento nessa propriedade o que indica um provável enrijecimento destes adesivos. Parte desse efeito devido ao tempo de armazenamento pode ter sido causado por possível absorção de umidade como insinuado pelos resultados de teor de voláteis. Embora essa relação não tenha sido verificada, essa possibilidade está indicada em literatura[4].

estabilizada (kit 4). Este aumento de voláteis credita-se ao fato da matriz de resinas epoxídicas e do reforço de poliamida serem intrinsecamente higroscópicos.

4.5 Solubilidade A Tabela 7 apresenta os resultados dos ensaios de solubilidade em solvente orgânico. Os resultados demonstram que o período de exposição ao ambiente tem grande influência na solubilidade da resina, sendo que o tempo de armazenamento não apresentou a mesma influência para essa propriedade. A diminuição da extração da resina pela presença de solvente indica uma redução da solubilidade desta após o tempo de exposição. Conforme aumenta o tempo de exposição do adesivo, a solubilidade da resina no solvente apresenta redução consistente. Lucas e co-autores explicam que esses aglomerados ocorrem pelo empacotamento das cadeias dificultando a difusão do solvente no material e que, quanto menor a distância intermolecular, menos a resina se solubilizará[22]. O aspecto dos adesivos, apresentado na Figura 4, se repetiu para todos os ciclos ensaiados, onde a figura representativa de 0 dia de exposição apresentou apenas fibras após a extração, e no outro extremo, com 12 dias de exposição, observa-se grande quantidade de resina não removida pelo solvente, evidenciando assim a perda de solubilidade. Nos períodos intermediários, notou-se diminuição constante da solubilidade, por meio da remoção parcial de resina em

4.4 Teor de voláteis A tabela 6 apresenta os resultados dos ensaios de teor de voláteis médios para os quatro ciclos de exposição e armazenamento. Apesar da metodologia indicar uma faixa de variação de temperatura larga (166,4°C a 177,6°C) o desvio-padrão dos resultados indica uma boa repetibilidade dos resultados encontrados, com exceção da amostra de 3 dias de exposição do kit 2. Conforme ocorre o aumento do período de exposição, os adesivos apresentaram um aumento no teor de voláteis ao longo dos 12 dias. Embora a causa não tenha sido investigada, estima-se que este aumento possa ter ocorrido devido à absorção de umidade. Mesmo os materiais estando embalados durante o período de armazenamento em freezer, o aumento de teor de voláteis é significativo entre os kits 1 a 3, após este período a absorção de umidade é

Tabela 5. Ensaio de aderência (nível). Exposição (dias) 0 3 5 8 10 12

Kit 1 IV IV III III III III

Kit 2 III III III III III II

Kit 3 III III III III II II

Kit 4 III III III II II II

Kit 1 0,68 ± 0,01 0,67 ± 0,02 0,75 ± 0,04 0,74 ± 0,01 1,10 ± 0,05 1,27 ± 0,03

Kit 2 1,35 ± 0,02 1,10 ± 0,37 1,26 ± 0,06 1,27 ± 0,02 1,30 ± 0,01 1,30 ± 0,01

Kit 3 2,01 ± 0,01 2,06 ± 0,03 2,04 ± 0,03 2,05 ± 0,02 2,12 ± 0,02 2,10 ± 0,02

Kit 4 2,00 ± 0,03 1,99 ± 0,01 2,03 ± 0,02 2,02 ± 0,02 2,03 ± 0,02 2,07 ± 0,03

Kit 1 95 ± 1 93* 87 ± 2 85 ± 1 75 ± 2 74 ± 2

Kit 2 96 ± 1 94 ± 1 87* 86 ± 1 72 ± 1 69 ± 1

Kit 3 93 ± 2 93 ± 3 88 ± 3 85 ± 2 81 ± 2 77 ± 2

Kit 4 94 ± 2 92 ± 2 88 ± 4 85 ±2 80 ± 6 73 ± 2

Tabela 6. Teor de voláteis (%). Exposição (dias) 0 3 5 8 10 12

Tabela 7. Solubilidade (%). Exposição (dias) 0 3 5 8 10 12

*Os resultados apresentaram desvio-padrão abaixo do número de algarismos significativo adotados para o ensaio.

Polímeros, 26(1), 92-100, 2016

97


Silva, A. C. T. N., Guilherme, F., Ferrari, V. M., & Ferrari, P. E.

Figura 4. Fotos do ensaio de solubilidade após extração da resina com solvente orgânico, o número indicado em cada foto corresponde ao período de envelhecimento de cada material em dias. Tabela 8. Dados obtidos da análise de Calorimetria Exploratória Diferencial (DSC) para o Kit nº 3. Exposição (dias) 0 3 5 8 10 12

Temperatura Onset (°C) 191 187 188 196 190 191

Temperatura de Pico (°C) 227 229 230 229 229 230

Entalpia ΔE (J/g) 258 301 294 239 227 239

quantidades cada vez menores. Estes resultados indicam que houve um aumento do número de reticulações no material termorrígido, que mesmo em baixas concentrações, dificulta a difusão do solvente dentro da massa polimérica.

4.6 Aspecto visual Observando-se as amostras a olho nú, antes dos ensaios, notou-se que, com 0 dia de exposição à temperatura ambiente, estas apresentavam uma superfície lisa, homogênea e brilhante. Para as amostras entre 3 e 5 dias de exposição, o adesivo já não era mais tão liso, apresentando também redução na homogeneidade e no brilho. Para os materiais entre 8 e 12 dias o material apresento-se rugoso, fosco e quebradiço.

4.7 Calorimetria Exploratória Diferencial (DSC) A partir dos resultados das curvas de DSC, nos diferentes tempos de exposição apresentados na Tabela 8, obtidos a partir das amostras do kit 3, foi possível observar que as variáveis Tonset, Tmáx e ΔE não apresentaram variações coerentes ou mudanças significativas na cura da resina relacionadas ao armazenamento e/ou envelhecimento à temperatura ambiente (Figura 5). Estes resultados sugerem que o possível aumento no número de reticulações da matriz polimérica à temperatura ambiente, não foi comprovado por este método de análise térmica, já que não houve alteração significativa da entalpia de cura ao longo do período de exposição, o que pode ser atribuído a uma taxa de cura muito baixa da matriz ou apenas pela falta sensibilidade do método. 98

Figura 5. Calorimétrica exploratória diferencial (DSC). Polímeros, 26(1), 92-100, 2016


Investigação do efeito do tempo de exposição à temperatura ambiente e ao tempo de estocagem de um filme adesivo estrutural de resina epoxídica

5. Conclusão De acordo com os resultados apresentados, independente do ciclo analisado, as amostras com 0 dia de envelhecimento, ou seja, adesivo com baixo tempo de exposição, apresentou elevada solubilidade da resina no solvente, elevada capacidade de fluir e altas propriedades de aderência dentro da escala adotada. Entre os materiais com 3 e 5 dias de envelhecimento, o adesivo apresentou uma pequena defasagem nos resultados quando comparados com os materiais de 0 dia. Neste caso, o material se encontra com o “tack” inferior, porém ainda apresenta boas características de aderência. Observa-se que no período de envelhecimento entre 5 e 8 dias o material sofreu um decréscimo na capacidade de fluxo e aderência, agravando-se conforme o aumento do tempo de exposição à temperatura ambiente. Nas propriedades de aderência e principalmente na solubilidade, observou-se visualmente que a resina aglomerou-se nas fibras de nylon e a matriz não flui mais de forma homogênea. Nos ensaios entre 10 e 12 dias de exposição à temperatura ambiente foi observado que os materiais apresentaram aparência fosca, pouca ou nenhuma aderência e a resina se aglomerou quase que totalmente nas fibras de nylon. O teste de solubilidade apresentou boa relação com o decréscimo de outras propriedades. Embora não se possa fazer uma correlação numérica direta, o teste de solubilidade mostrou-se sensível à transformação do material demonstrando que, por se tratar de um parâmetro de avaliação rápido, simples e barato, sem a necessidade de equipamentos auxiliares, é um instrumento potencial de avaliação muito forte para uso no meio produtivo. Finalmente, conclui-se que o a exposição à temperatura ambiente tem efeito deletério generalizado nas propriedades avaliadas, mas quando armazenados à baixa temperatura, mesmo por tempo exponencialmente superior aos avaliados em temperatura ambiente, as propriedades sofreram pouca ou alteração não detectável. Assim, a exposição prolongada de materiais não curados às condições de temperatura e umidade acima das recomendadas pelo fabricante contribuem na redução do desempenho e podem se tornar causas potenciais de falha prematura da ligação das superfícies unidas por adesivos. Para o futuro, sugere-se uma análise mais detalhada sobre o assunto, inclusive com outros sistemas de resinas para avaliar a dependência dessas propriedades com relação à formulação. Também é interessante aumentar o período de armazenamento para verificar a transformação nessa condição e tentar correlacionar com dados de envelhecimento ao ambiente de forma a utilizar este último como ensaio acelerado. Sugere-se ainda a complementação do estudo com outras técnicas para avaliar o grau de cura da matriz polimérica e ensaios mecânicos para avaliação do impacto das transformações detectadas nas propriedades finais de colagem.

6. Referências 1. American Society for Testing and Materials – ASTM. (2011). ASTM D907: standard terminology of adhesives. West Conshohocken: ASTM. Polímeros, 26(1), 92-100, 2016

2. Sharpe, L. H. (1990). Section1. Fundamentals of adhesives and sealants technology. In ASM. Engineered materials handbook – adhesives and sealants (Vol. 3, pp. 1-72). Materials Park: ASM International. 3. Hexcel. (2003). Handbook Redux® Bonding Technology. Stamford: Hexcel Corporation. Recuperado em 10 de setembro de 2013, de http://www.hexcel.com/Resources/DataSheets/ Brochure-Data-Sheets/Adhesive_Bonding_Technology.pdf 4. Petrie, E. M. (2007). Handbook of adhesives and selants (2. ed). New York: The McGraw-Hill Companies. 5. Campbell, F. C. (2006). Manufacturing technology for aerospace structural materials. Amsterdam: Elsevier Science. 6. Schwartz, M. M. (1997). Composite materials: properties, nondestructive testing, and repair. New Jersey: Prentice-Hall Inc. 7. Rezende, M. C., Botelho, E. C., Paiva, J. M. F., & Costa, M. L. (2003). Avaliação Térmica e reológica do ciclo de cura do pré-impregnado de carbono epóxi. Polímeros: Ciência e Tecnologia, 13(3), 188-197. http://dx.doi.org/10.1590/S010414282003000300009. 8. Callister, W.D., Jr., & Rethwisch, D. G. (1991). Materials science and engineering an introduction. New York: John Wiley & Sons. 9. Fraga, F., Castro-Díaz, C., Rodrıguez-Nuñez, E., & MartínezAgeitos, J. M. (2003). Physical aging for an epoxy network diglycidyl ether of bisphenol A/m-xylylenediamine. Polymer, 44(19), 5779-5784. http://dx.doi.org/10.1016/S00323861(03)00624-4. 10. Riegel, I. C., Freitas, L. L., & Samios, D. (1999). Envelhecimento físico de sistemas DGEBA/DDM investigado por análise térmica (DSC/DMA). Polímeros: Ciência e Tecnologia, 9(3), 58-64. http://dx.doi.org/10.1590/S0104-14281999000300011. 11. Kim, D., Centea, T., & Nutt, R. (2014). Out-time effects on cure kinetics and viscosity for an out-of-autoclave (OOA) prepreg: Modelling and monitoring. Composites Science and Technology, 100, 63-69. http://dx.doi.org/10.1016/j. compscitech.2014.05.027. 12. Sales, R. C. M., & Dibbern-Brunelli, D. (2008). Efeito higrotérmico em prepregs de fibra de vidro/epóxi por espectroscopia de luminescência. Polímeros: Ciência e Tecnologia, 18(1), 52-56. http://dx.doi.org/10.1590/S0104-14282008000100011. 13. Yagoubi, J. E., Lubineau, G., Traidia, A., & Verdu, J. (2015). Monitoring and simulations of hydrolysis in epoxy matrix composites during hygrothermal aging. Composites. Part A, Applied Science and Manufacturing, 68, 184-192. http://dx.doi. org/10.1016/j.compositesa.2014.10.002. 14. Ray, B. C. (2006). Temperature effect during humid ageing on interfaces of glass and carbon fibers reinforced epoxy composites. Journal of Colloid and Interface Science, 298(1), 111-117. http://dx.doi.org/10.1016/j.jcis.2005.12.023. PMid:16386268. 15. Sales, R. C. M., Diniz, M. F., Dutra, R. C. L., Thim, G. P., & Dibbern-Brunelli, D. (2010). Thermal curing of glass-epoxy prepregs by luminescence spectroscopy. Journal of Applied Polymer Science, 117(2), 664-671. http://dx.doi.org/10.1002/ app.31953. 16. Twisk, J. V., & Aker, S. C. (1990). Storing adhesives and sealant materials. In ASM. Engineered materials handbook – adhesives and sealants (Vol. 3, pp. 683-686). Materials Park: ASM International. 17. American Society for Testing and Materials – ASTM. (1999). ASTM D3532/D3532M: gel time of carbon fiber-epoxy prepreg. West Conshohocken: ASTM. 18. American Society for Testing and Materials – ASTM. (1997). ASTM D3530/D3530-M: volatiles content of composite material prepreg.West Conshohocken: ASTM. 99


Silva, A. C. T. N., Guilherme, F., Ferrari, V. M., & Ferrari, P. E. 19. American Society for Testing and Materials – ASTM. (2011). ASTM D3531/D3531-M: resin flow of carbon fiber-epoxy prepreg.West Conshohocken: ASTM. 20. National Center for Advanced Materials Performance – NCAMP. (2015). NMS 5320: material specification. Wichita: National Institute for Aviation Research. 21. Rhodia Solvay Group. (2013). Rhodiasolv PolarClean. Solvay. Recuperado em 09 de setembro de 2013, de http://www.

100

rhodia.com.br/pt/mercados-e-produtos/catalogo-de-produtos/ Rhodiasolv-PolarClean.html 22. Lucas, E. F., Soares, B. G., & Monteiro, E. E. C. (2001). Caracterização de polímeros: determinação de peso molecular e análise térmica. Rio de Janeiro: E-papers. Enviado: Set. 17, 2014 Reenviado: Ago. 28, 2015 Aceito: Set. 08, 2015

Polímeros, 26(1), 92-100, 2016


DRIVING INNOVATION FORWARD Polímeros

New materials like plastics are advancing vehicle technologies. But getting the best results is a big challenge. To assist, SABIC offers industry-leading expertise in designing with a wide range of engineered materials. Across the entire vehicle. Because no matter what obstacles may hold our customers back, we’re there with ‘Chemistry that MattersTM’ to help them drive forward.

SABIC.com

VOLUME XXVI - N° 1 - JAN/FEV/MAR - 2016

© 2016 Copyright SABIC. All rights reserved.

Polímeros Ciência e Tecnologia, vol.26, n.1, 2016  

The journal Polímeros: Ciência e Tecnologia is a quarterly publication of the Brazilian Polymer Association (Associação Brasileira de Políme...

Read more
Read more
Similar to
Popular now
Just for you