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P o l í m e r o s - N º 2 - V o l u m e X X V I - A b r / m A I o / J u N - 2 0 1 6 - I s s N 0 1 0 4 - 1 4 2 8 - I s s N 1 6 7 8 - 5 1 6 9

I n d e x a d a : “ C h e m I C a l a b s t r a C t s ” — “ r a P r a 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 C I e n C e a n d t e C h n I C 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 I C a s d e a m e r I C 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 C I e n C e ”

P o l í m e r o s

Pr e s I d e N t e d o Co N s e l h o ed I t o r I A l

Marco-Aurelio De Paoli (UNICAMP/IQ)

me m b r o s d o Co N s e l h o ed I t o r I A l

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)

Polímeros / Associação Brasileira de Polímeros. vol. 1, nº 1 (1991) -.- São Carlos: ABPol, 1991-

Trimestralv. 26, nº 2 (Abr./Maio/Jun. 2016)ISSN 0104-1428ISSN 1678-5169 (versão eletrônica)

1. Polímeros. l. Associação Brasileira de Polímeros.

“Polímeros” é uma publicação daAssociação Brasileira de Polímeros

Rua São Paulo, nº 99413560-340 - São Carlos, SP, Brasil

Fone/Fax: (16) 3374-3949

e-mails: abpol@abpol.org.br / revista@abpol.org.brhttp://www.abpol.org.br

Data de publicação: Junho de 2016

Versão eletrônica disponível no site:www.scielo.br

Apoio:

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

Pr o d u ç ã o e As s e s s o r I A ed I t o r I A l

www.editoracubo.com.br

Co m I t ê ed I t o r I A l

Sebastião V. Canevarolo Jr. – Editor

me m b r o s d o Co m I t ê ed I t o r I A l

Adhemar C. Ruvolo FilhoAlain Dufresne

Bluma G. SoaresCésar Liberato Petzhold

João B. P. SoaresJosé António C. Gomes Covas

José Carlos C. S. PintoRegina Célia R. Nunes

Richard G. WeissRodrigo Lambert Oréfice

( V e r s ã o e l e t r ô N I C A )

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se ç ã o ed I t o r I A lInformes & Notícias ...........................................................................................................................................................................E3Associados ..........................................................................................................................................................................................E5Calendário de Eventos .......................................................................................................................................................................E7

se ç ã o té C N I C A

Thermo stabilisation of poly (butylene adipate-co-terephthalate)Rodrigo Paulino Chaves and Guilhermino José Macêdo Fechine ................................................................................................................. 102

Layered double hydroxides as fillers in poly(l-lactide) nanocomposites, obtained by in situ bulk polymerizationTelma Nogueira, Núria Gonçalves, Rodrigo Botan, Fernando Wypych and Liliane Lona ............................................................................. 106

Curaua fiber reinforced high-density polyethylene composites: effect of impact modifier and fiber loadingJaqueline Albano de Morais, Renan Gadioli and Marco-Aurelio De Paoli ................................................................................................... 115

Effect of compatibilization in situ on PA/SEBS blendsAnna Paula Azevedo de Carvalho and Alex da Silva Sirqueira ...................................................................................................................... 123

Structure-flammability relationship study of phosphoester dimers by MLR and PLSLuminita Crisan, Smaranda Iliescu and Simona Funar-Timofei .................................................................................................................... 129

Influence of PLGA and PLGA-PEG on the dissolution profile of oxaliplatinEmiliane Daher Pereira, Renata Cerruti, Edson Fernandes, Luis Peña, Vivian Saez, José Carlos Pinto, José Angel Ramón, Geiza Esperandio Oliveira and Fernando Gomes de Souza Júnior ............................................................................................................... 137

Biopolymer production using fungus Mucor racemosus Fresenius and glycerol as substrateThaíssa Rodrigues Araújo, Carmen Lúcia de Oliveira Petkowicz, Vicelma Luiz Cardoso, Ubirajara Coutinho Filho and Patrícia Angélica Vieira .................................................................................................................................................................................. 144

Influence of nucleating agent on the crystallization kinetics and morphology of polypropyleneAdriane Gomes Simanke, Ana Paula de Azeredo, Cristóvão de Lemos and Raquel Santos Mauler .............................................................. 152

Biodegradation of additive PHBV/PP-co-PE films buried in soilBarbara Rani-Borges, Adriano Uemura Faria, Adriana de Campos, Suely Patricia Costa Gonçalves and Sandra Mara Martins-Franchetti ...... 161

The effect of andiroba oil and chitosan concentration on the physical properties of chitosan emulsion filmVanessa Tiemi Kimura, Cintia Satiyo Miyasato, Bianca Pereira Genesi, Patrícia Santos Lopes, Cristiana Maria Pedroso Yoshida and Classius Ferreira da Silva ........................................................................................................................................................................ 168

Preparação e caracterização de poliuretanos contendo diferentes quantidades de óleo de baruElizabeth Luiza de Almeida, Gilberto Alessandre Soares Goulart, Salvador Claro Neto, Gilberto Orivaldo Chierice e Adriano Buzutti de Siqueira ............................................................................................................................................................................ 176

Caracterização de pinos da blenda poli(L-co-D,L ácido láctico)/poli(caprolactona triol) (PLDLA/PCL-T) e análise das propriedade mecânicas dos pinos durante degradação in vitroMarcia Adriana Tomaz Duarte, Ariana Cristina Motta e Eliana Aparecida de Rezende Duek ..................................................................... 185

Capa: SEM micrographs of the cryogenic fracture surface of injection molded composites: (a) HDPE/20EVA20CF showing fiber-matrix interaction and co-continuous EVA phase, (b) HDPE/30EVA20CF showing the homogeneous distribution of the fibers,

(c) HDPE/40EVA20CF highlighting fiber diameter and (d) HDPE/40EVA20CF composite after extraction with acetone with higher magnification.

Elaboração artística Editora Cubo.

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Using terahertz laser, scientists change the macromolecular conformation of a polymer

Scientists from the RIKEN Center for Advanced Photonics (RAP) have, for the first time, successfully used a terahertz laser to induce permanent changes in the conformation of a polymer, giving it an increased pattern of crystallization. Conformational changes are very important for macromolecular science because they can change the characteristics of a material and, in the case of proteins, can make it either possible or impossible for them to perform a certain biological function. The work, done in collaboration with Osaka University, was published in Scientific Reports.

According to Hiromichi Hoshina of RAP, “Terahertz lasers offer promise as a way to modify materials, because they resonate at a frequency close to the oscillations of the hydrogen bonds that bind polymers into certain conformations, but are much lower in energy than the covalent bonds that make up the molecular structure of the polymers. As a result, they could offer a ‘soft’ way to change the conformation without inducing chemical changes.”

One of the difficulties, however, of using terahertz wave irradiation to induce changes is that the materials tend to revert very quickly to the state of thermal equilibrium states. To overcome this challenge, the group decided to perform experiments on a polymer undergoing solvent casting crystallization—a process through which the conformation is fixed. This allowed them to effectively “fix” the results of the work and detect any changes.

The experiment was successful. When the group irradiated a polymer—a poly(3-hydroxybutylate)/chloroform solution—with terahertz radiation with a peak power of 40 megawatt/cm2, using a terahertz free electron laser FEL - developed by the Institute of Scientific and Industrial Research at Osaka University, they found that the crystallization of the material was increased by 20%.

“We were happy with these results, but we were also surprised by what we saw,” continues Hoshina.”The researchers were intrigued, however, by the fact that the peak power used in this study was quite lower than previous reports using NIR and visible lasers. They considered that the crystallization might have been caused by changes in temperature, but measured it and found that the difference between regions was less than 1 degree Celsius, much too small a difference to account for the difference. They also considered that the terahertz waves might have directed caused increased vibrations between the molecules but did not find any significant correlations with the wavelength—something that should have happened if the effect was due to differences in resonance.

According to Hoshina, “We have, for the first time, shown that terahertz waves can effectively induce a rearrangement of the molecules in polymer macromolecules. The exact mechanism through which

this happens remains a mystery, though we speculate that it might be related to the generation of shockwaves in the material, and we plan future work to find out exactly what is special about these terahertz waves, which have often been called the ‘unexplored frontier of the electromagnetic spectrum’.”

“We are excited by this work,” he continues, “as this could give us a new tool for controlling the structure of ‘fragile’ molecules and allowing us to discover new functional materials.”

Source: Phys.org

US BIOPOLYMERS

US demand for natural polymers (biopolymers) is forecast to increase at a rate of 4.3% annually, reaching USD 5.1 bn in 2020, or 862,000 t, in consumption. Growth will be driven by demand for natural ingredients in the food and beverage industry and the medical market, supporting use of cellulose ethers and starch and fermentation polymers, says US market researcher Freedonia (Cleveland, Ohio; www.freedoniagroup.com) in its recently published study, “Natural polymers”.

Cellulose ethers are the largest product type, representing a third of the market, and methyl cellulose is expected to remain market leader through the forecast period. However, demand for hydroxyethyl cellulose (HEC), microcrystalline cellulose (MCC), and carboxymethyl cellulose (CMC) will also be significant. Methyl cellulose’s largest application is in construction, where it has a variety of uses, including plastering, flooring, grouting, mortaring, tile adhesion and stucco.

Starch and fermentation products will advance the most rapidly of all natural polymers through 2020, with PLA accounting for much of the growth. As production capacities have increased, PLA is one of the few natural polymers that has exhibited declining prices. This has been a great benefit for the packaging segment, which represents by far the leading market for PLA resins, the study says.

Freedonia analyst Larry Catsonis said: “Hyaluronic acid will also support starch and fermentation products gains, as it continues to be used to relieve joint pain in orthopedic injections and will also gain market share from collagen as both a dermal and topical tissue filler in cosmetic applications.”

The food and beverage industry will relinquish its spot as the largest market for natural polymers to the medical industry in the forecast period, Freedonia said. Growth in medical applications will be driven by strong demand from collagen in wound care as well as cellulose ethers used in pharmaceuticals.

Additionally, declining oil and gas production, mirroring price declines into 2015, led to less demand for oilfield natural polymers, especially guar gum, which is used as a carrier for placing sand into fractures and as a top-hole drilling fluid.

Source: Plasteurope.com

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Anúncio - AX Plásticos - 190x142 mm.pdf 1 01/06/16 14:19

Associados da ABPol

Patrocinadores

Instituições

UFSCar/ Departamento de Engenharia de Materiais, SPSENAI/ Serviço Nacional de Aprendizagem Industrial Mario Amato, SPUFRN/ Universidade Federal do Rio Grande do Norte, RN

Polímeros, 26(2), 2016 E5

Associados da ABPol

Coletivos

A. Schulman Plásticos do Brasil Ltda.Aditive Plásticos Ltda.Avamplas – Polímeros da Amazônia Ltda.CBE – Grupo UnigelColorfix Itamaster Indústria de Masterbatches Ltda.Cromex S/ACytec 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/AInstituto de Aeronáutica e Espaço/AQIJaguar Ind. e Com. de Plásticos LtdaJohnson & 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

E6 Polímeros, 26(2), 2016

August

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

3rd Brazilian Conference on Composite Materials (BCCM-3)Date: 28-31 August 2016Local: Gramado - RSWebsite: http://www.bccm.com.br

BiPoCo 2016 - 3rd International Conference on Bio-based Polymers and CompositesDate: August 28 - September 1, 2016Local: Szeged - HungaryWebsite: http://bipoco2016.hu/

September

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

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

International Conference on Advanced Energy Materials (AEM2016)Date: 12–14 September 2016Local: Guildford (Surrey) - United KingdomWebsite: http://www.aem2016.com/

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

PLASTEC MinneapolisDate: 21–22 September 2016Local: Minnesota - USAWebsite: http://plastecminn.plasticstoday.com/

Organic SemiconductorsDate: 22–25 September 2016Local: Dubrovnik - CroatiaWebsite: http://www.zingconferences.com/conferences/organic-semiconductors/

4th Symposium on Innovative Polymers for Controlled Delivery (SIPCD 2016)Date: 23–26 September 2016Local: Suzhou - ChinaWebsite: http://www.sipcd.com/

Polyurethanes Technical ConferenceDate: 26–28 September 2016Local: Maryland - USAWebsite: www.polyurethane.americanchemistry.com

Colombiaplast 2016Date: 26–30 September 2016Local: Bogotá - ColombiaWebsite: http://www.colombiaplast.com/

Conductive Plastics - 2016Date: 26–30 September 2016Local: Pennsylvania - USAWebsite: http://www.amiplastics.com/events/event?Code=C742

October

Polymeric Implants & Catheters in Medical DevicesDate: 4–6 October 2016Local: Las Vegas - USAWebsite: http://www.mediplastconference.com/

IUPAC International Conference on Advanced Polymeric MaterialsDate: 4–7 October 2016Local: Jeju - South KoreaWebsite: http://www.psk40.org/

China International Exhibition on Plastics and Rubber Injection Moulding Industry (CIM) 2016Date: 13–15 October 2016Local: Tianjin - ChinaWebsite: http://www.cimexpo.cn/

November

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

36th Australasian Polymer SymposiumDate: 20–23 November 2016Local: Lorne - AustraliaWebsite: http://www.36aps.org.au/

3º. Encontro Nordeste de Ciência e Tecnologia de PolímerosDate: 28–30 November 2016Local: Fortaleza - CearáWebsite: http://www.abpol.org.br

11th European Bioplastics ConferenceDate: 29–30 November 2016Local: Berlin - Germany Website: www.european-bioplastics.org

Composites Europe 2016Date: November 29 - December 1, 2016Local: Messe Düsseldorf, Germany Website: www.composites-europe.com

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

December

Polymers in Flooring – 2016Date: 6-7 December 2016Local: Berlin - GermanyWebsite: http://www.amiplastics.com/events/event?Code=C769

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

Polímeros, 26(2), 2016 E7

http://dx.doi.org/10.1590/0104-1428.2196

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Thermo stabilisation of poly (butylene adipate-co-terephthalate)

Rodrigo Paulino Chaves1 and Guilhermino José Macêdo Fechine2*

1Materials Engineering Department, Engineering School, Universidade Presbiteriana Mackenzie, São Paulo, SP, Brazil

2Graphene and Nano-materials Research Center – MackGraphe, Universidade Presbiteriana Mackenzie, São Paulo, SP, Brazil

*guilherminojmf@mackenzie.br

Sbstract

Poly (butylene adipate-co-terephthalate) - PBAT is a synthetic biodegradable polymer commonly used for plastic film production from neat polymer or nanocomposites. The PBAT is submitted to high temperatures and shear rate during its processing. In the present study, the thermo stabilisation of PBAT by the addition of two types of stabilisers was studied using a torque rheometer at 60 rpm and two levels of temperature. The stabilisers were used as master batches with a percentage of 10% by weight of additive in the PBAT. Molecular weight, torque values after 10 minutes of mixing, and absorbance at 400 nm were used to evaluate the process of stabilisation. The primary and secondary antioxidant used here had a positive effect on both processing temperatures, 180 and 200 °C. The best results indicate that the primary antioxidant could be used alone to protect PBAT against thermodegradation reactions.

Keywords: biodegradable polymer, poly (butylene adipate-co-terephthalate), thermo stabilisation.

1. Introduction

The main applications of biodegradable polymers do not require high mechanical strength like packaging, disposable non-woven, sanitary products, consumables and agricultural tools[1]. Biodegradable polymers still face some problems in their use due to their low performance when subjected to applications that need high strength, whether chemical, physico-chemical and/or mechanical. However, advancements in research for better mechanical properties in biodegradable polymers[2,3] and environmental problems like pollution, largely influenced by conventional polymers[4,5] led to the enhancement of biodegradable polymer production. The greatest concern about biodegradable polymers is the time of the biodegradation and bioassimilation of them during the degradation process induced by microorganisms.[6,7] The biodegradation of some polymers is governed by the attack of the micro-organisms at ester linkage that enables a rapid fragmentation. This same linkage is very sensitive to the degradative process caused by high temperature and shear rates as well as hydrolysis due to the presence of moisture. The exposure to these degradation conditions could occur mainly during processing (extrusion and injection moulding). Few researchers have presented studies about thermostabilisation of the biodegradable polymers[8]. Poly (butylene adipate-co-terephthalate) - PBAT is a synthetic biodegradable copolymer, specifically a copolyester of adipic acid, 1, 4-butanediol and dimethyl terephthalate. It is commonly used for plastic film production[9]. The polymer processing to produce films or to obtain nanocomposites based on PBAT is carried out at high temperatures and shear rate during its processing. Al-Itry et al.[10] proposed the degradation mechanism of PBAT during processing based on hydrolysis of ester linkage, main-chain scissions and β-C-H hydrogen transfer. The choice of stabilisers is

very important to keep the physical properties of PBAT after processing and it is dependent of action mechanism of the stabilisers. Here, the objective of this study is to evaluate the thermo-mechanical stabilisation of PBAT through the use of antioxidants (primary and secondary), using a torque rheometer as a processing and analysis tool. Torque monitoring after 10 minutes, analysis of molecular weight measurements by size exclusion chromatography (SEC) and UV/Vis were used to assess degradation and stabilisation of PBAT. The idea includes not only the evaluation of thermostabilisation of PBAT but also simplifies the method of this analysis.

2. Experimental

In this work, commercial PBAT was used. The stabilisers used in this work were Irganox 1010 (primary antioxidant - P, 0.4 w/w%) and Irgafos 168 (secondary antioxidant - S, 0.5 w/w%). As the concentrations of additives are very low, it was preferred to prepare master batches with a percentage of 10% by weight of additive in the PBAT. After that, a fresh mixture with pure PBAT and master batch was made to achieve the desired concentrations. The preparation of the concentrates was carried out in a mechanical mixer and then they were cut in a knife mill to obtain smaller sizes. The thermo-mechanical degradation was performed on a torque rheometer at 60 rpm for 10 min. Two levels of temperature inside the rheometer were used, 180 and 200 °C. PBAT and PBAT/master batch were dried in an oven for 1 hour at 70 °C before processing. The evaluation of the stabilisation process was done by analysing the torque value after 10 minutes, Size Exclusion Chromatography (SEC)

Thermo stabilisation of poly (butylene adipate-co-terephthalate)

Polímeros, 26(2), 102-105, 2016 103

and UV/Vis spectroscopy. SEC analyses were conducted in a Viscoteck with a series of columns at 40 °C and with a refractive index detector. Specimens were dissolved in THF and the filtered solution was injected into the equipment. The solvent flow rate (THF) was 1 mL/ min and the columns were calibrated with narrow molecular weight PS. The UV-visible spectroscopy analysis was done with the samples in chloroform solution (0.01 g/cm3) using a Shimadzu 1501 equipment.

3. Results and Discussion

Table 1 presents values of nM , wM , polydispersity (PD), torque after 10 minutes, and absorbance at 400 nm for neat PBAT and PBAT + additives processed at the rheometer for 10 minutes under 180 and 200 °C. Molecular weight data in Table 1 indicates that temperature has a great influence on the thermodegradation of PBAT since there is a higher drop of molecular weight values for samples without stabilisers processed at 200 °C when compared with those processed at 180 °C. At the same time, the value of polydispersity has increased. These results indicated that the process involved in the thermodegradation of PBAT is governed by the scission chain reactions[11,12]. This hypothesis is confirmed by the rheometric data since the values of torque at 10 min has decreased significantly at 200 °C. In the case of the additives at 180 °C, there is no significant difference between the action of primary and secondary antioxidants taking into account only molecular weight and torque data. Values of molecular weight remained almost unchanged for 180 °C compared with neat PBAT and torque values at 10 min were quite similar for the three compositions. It is a good indication of thermostabilisation of PBAT by the two additives used here. However, the expected synergetic effect between primary and secondary antioxidant did not occur clearly. Molecular weight data for samples with additives processed at 200 °C under process conditions used here show that the additives also have a positive effect. There is no trend with regard to polydispersity; it is probable that scissions and crosslinking reactions competed during the thermodegradation process of PBAT. Torque values at 10 minutes are lower than samples processed at 180 oC, therefore, higher than samples without additives. Higher

temperatures led to a strong thermodegradation of PBAT even in a presence of additives. Absorbance at 400nm of the polymer solutions are used to describe two types of phenomena, an increase of chromophore groups or high dispersion of the light due to insoluble fragments of polymer from crosslinking reactions[13]. In both temperatures, the presence of additives decreased the absorbance at 400 nm indicating the positive action of the stabilisers against of thermodegradation reactions. The highest decrease of Abs400nm is for the sample processing at 180 °C in the presence of secondary stabiliser. Probably, the additives could prevent the crosslinking reactions at this temperature but not the chain scissions since the drop of molecular weight is quite the same for the samples processing with primary stabiliser alone or in combination with secondary one.

Figure 1 shows rheometer torque curves of PBAT processed at 180 and 200 °C without and in the presence of stabilisers. As can be seen, all compositions with stabilisers at both temperatures led to torque values above PBAT without stabilisers. The action of stabilisers was stronger for lower temperature (180 °C) and the primary oxidant alone was more efficient than the secondary one and the combination of primary and secondary. The primary stabilisers act directly on the deactivation of free radicals and the secondary ones act on the deactivation of free radicals or hydroperoxide decomposition[14]. It is an indication that thermostabilisation of PBAT is strongly driven by the deactivation of the free radicals when compared with the decomposition of hydroperoxides.

Figure 2 shows molecular weight distribution curves of neat PBAT and PBAT processed at 180 and 200 °C without stabilisers and with stabilisers. In the case of PBAT processed at both temperatures, clearly, it is verified that the molecular weight curve is displaced strongest to lower molecular weight when the polymer is processed with no stabiliser. However, the best stabiliser effect is obtained by the presence of the primary antioxidant alone for both temperature conditions. These results are in agreement with torque and UV/Vis results. It could be an indication that deactivation of free radicals reactions are more important to PBAT than hydroperoxide decomposition ones or the mechanism of degradation of PBAT is not governed by generation of hydroperoxides groups.

Table 1. Molecular weight, torque and UV/Vis data of PBAT samples processed without and with stabilisers at two different temperatures.

Sample/Temperature nM(g/mol)

wM(g/mol)

PDT10min

(N.m)Abs400nm

(u.a.)Neat PBAT 40.600 84.400 2.08 - 0.00829PBAT/180 °C 36.450 76.950 2.11 1.8 0.03297PBAT/200 °C 33.250 70.200 2.11 1.5 0.03928PBAT + P/180 °C 40.500 80.000 1.97 4.5 0.02516PBAT + S/180 °C 39.750 80.900 2.03 4.1 0.01699PBAT + PS/180 °C 39.700 79.100 1.99 4.0 0.02539PBAT + P/200 °C 40.100 78.300 1.95 2.6 0.02216PBAT + S/200 °C 36.850 74.000 2.01 2.5 0.02516PBAT + PS/200 °C 38.300 74.700 1.95 2.0 0.02321

Chaves, R. P., & Fechine, G. J. M.

Polímeros, 26(2), 102-105, 2016104

Figure 1. Rheometer torque curves of PBAT processed at 180 (a) and 200 °C (b) without stabilisers and with stabilisers P = primary antioxidant S = secondary oxidant, PS = P+S. Detailed curves around 10 minutes for 180 (c) and 200 °C (d).

Figure 2. Molecular weight distribution curves of neat PBAT and PBAT processed at 180 (a) and 200 °C (b) without stabilisers and with stabilisers P = primary antioxidant S = secondary oxidant, PS = P+S.

4. Conclusions

In this work, the thermo stabilisation of PBAT was studied using a torque rheometer and two types of stabilisers (primary and secondary antioxidant). Molecular weight, torque values after 10 minutes of mixing, and absorbance at 400 nm were

used to evaluate the process of stabilisation. The results show that these two types of stabilisers worked very well, both alone and together, however, with the concentrations used here the primary antioxidant is the best choice to stabilise PBAT during processing. However, a better control of the drying of the samples must be performed to avoid

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Polímeros, 26(2), 102-105, 2016 105

the hydrolysis reactions which can not be prevented by the action of stabilisers. Here, it is very important to highlight that nothing could be predicted concerning prolonged use of the stabilisers since the data were acquired on a ten-minute range. The torque rheometer technique proved to be able to generate results in a simple way and with indicators for the best choice of stabiliser.

5. Acknowledgements

The authors are grateful to FAPESP (Process 2010/07651-9) and CNPq (Process 304902/2012-8) Brazilian funding agencies for the financial support.

6. References

1. Bastioli, C. (2005). Handbook of biodegradable polymers. United Kingdom: Rapra Technology.

2. Lee, S. M., Cho, D., Park, W. H., Lee, S. G., Han, S. O., & Drzal, L. T. (2005). Novel silk/poly(butylene succinate) biocomposites: the effect of short fibre content on their mechanical and thermal properties. Composites Science and Technology, 65(3-4), 647-657. http://dx.doi.org/10.1016/j.compscitech.2004.09.023.

3. Bordes, P., Pollet, E., & Avérous, L. (2009). Nano-biocomposites: Biodegradable polyester/nanoclay systems. Progress in Polymer Science, 34(2), 125-155. http://dx.doi.org/10.1016/j.progpolymsci.2008.10.002.

4. Baeyens, J., Brems, A., & Dewil, R. (2010). Recovery and recycling of post-consumer waste materials. Part 2. Target wastes (glass beverage bottles, plastics, scrap metal and steel cans, end-of-life tyres, batteries and household hazardous waste). International Journal of Sustainable Engineering, 3(4), 232-245. http://dx.doi.org/10.1080/19397038.2010.507885.

5. Coelho, T. M., Castro, R., & Gobbo, J. A., Jr (2011). PET containers in Brazil: Opportunities and challenges of a logistics model for post-consumer waste recycling. Resources, Conservation and Recycling, 55(3), 291-299. http://dx.doi.org/10.1016/j.resconrec.2010.10.010.

6. Chandra, R., & Rustgi, R. (1998). Biodegradable polymers. Progress in Polymer Science, 23(7), 1273-1335. http://dx.doi.org/10.1016/S0079-6700(97)00039-7.

7. Luckachan, G. E., & Pillai, C. K. S. (2011). Biodegradable polymers: a review on recent trends and emerging perspectives. Journal of Polymers and the Environment, 19(3), 637-676. http://dx.doi.org/10.1007/s10924-011-0317-1.

8. Amorin, N. S. Q. S., Rosa, G., Alves, J. F., Gonçalves, S. P. C., Franchetti, S. M. M., & Fechine, G. J. M. (2014). Study of thermodegradation and thermostabilization of poly(lactide acid) using subsequent extrusion cycles. Journal of Applied Polymer Science, 131(6), 1-8. http://dx.doi.org/10.1002/app.40023.

9. Bilck, A. P., Grossmann, M. V. E., & Yamashita, F. (2010). Biodegradable mulch films for strawberry production. Polymer Testing, 29(4), 471-476. http://dx.doi.org/10.1016/j.polymertesting.2010.02.007.

10. Al-Ltry, R., Lamnawara, K., & Maazouz, A. (2012). Improvement of thermal stability, rheological and mechanical properties of PLA, PBAT and their blends by reactive extrusion with functionalized epoxy. Polymer Degradation & Stability, 97(10), 1898-1914. http://dx.doi.org/10.1016/j.polymdegradstab.2012.06.028.

11. Rabello, M. S., & White, J. R. (1997). Fotodegradação do polipropileno: um processo essencialmente heterogêneo. Polímeros: Ciência e Tecnologia, 7(2), 47-57. http://dx.doi.org/10.1590/S0104-14281997000200007.

12. Cáceres, C. A., & Canevarolo, S. V. (2008). Cisão de cadeia na degradação termo-mecânica do poliestireno sob múltiplas extrusões. Polímeros: Ciência e Tecnologia, 18(4), 348-352. http://dx.doi.org/10.1590/S0104-14282008000400015.

13. Timóteo, G. A. V., Fechine, G. J. M., & Rabello, M. S. (2007). Stress Cracking and Photodegradation: The Combination of Two Major Causes of HIPS Failure. Macromolecular Symposia, 258(1), 162-169. http://dx.doi.org/10.1002/masy.200751218.

14. Rabello, M. S., & De Paoli, M. (2013). Aditivação de termoplásticos, São Paulo: Artliber Editora Ltda.

Received: May 20, 2015 Revised: Jan. 11, 2016

Accepted: Feb. 15, 2016

http://dx.doi.org/10.1590/0104-1428.2282

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Layered double hydroxides as fillers in poly(l-lactide) nanocomposites, obtained by in situ bulk polymerization

Telma Nogueira1*, Núria Gonçalves1, Rodrigo Botan1, Fernando Wypych2 and Liliane Lona1

1Laboratory of Analysis, Simulation and Synthesis of Chemical Process – LASSPQ, Department of Materials Engineering and Bioprocesses, School of Chemical Engineering,

Universidade Estadual de Campinas – UNICAMP, Campinas, SP, Brazil2Advanced Materials Chemistry Laboratory, Chemistry Department, Universidade Federal

do Paraná – UFPR, Curitiba, PR, Brazil*telma.nogueira@gmail.com

Sbstract

In this study in situ bulk polymerization of L-lactide filled with layered double hydroxides (LDH) was investigated. Four different LDHs intercalated with two different organic anions (salicylate and sebacate) were synthesized and characterized. After characterization, these synthetic layered compounds were used as fillers in poly(L-lactide) (PLLA) nanocomposites with two different fillers’s loadings (1 wt% and 2 wt%). PLLA and PLLA nanocomposites were evaluated by X-ray diffraction (XRD), Fourier transform infrared spectroscopy (FTIR), ultraviolet and visible spectroscopy, thermogravimetric analysis (TGA), dynamical mechanical analysis (DMA), flexural testing and differential scanning calorimetry (DSC). The results demonstrated that, compared to PLLA, the nanocomposite containing 1 wt% of Zn/Al salicylate transmitted less UVA and UVB light, while keeping a similar transparency in the visible region. Thermogravimetric analysis revealed that the nanocomposite with 1 wt% of Zn/Al salicylate exhibited the highest thermal stability. In general the flexural and dynamical mechanical properties were reduced in compassion to neat PLLA. DSC results, demonstrated that, compared to PLLA, all the nanocomposites exhibited lower glass transition temperature and melting temperature values.

Keywords: layered double hydroxide, polylactide, polymeric nanocomposites, in-situ polymerization.

1. Introduction

Nowadays, the reduction of the environmental impact caused by non-biodegradable polymers is a growing interest, especially when they are applied in the production of disposable items, such as packaging. For such applications nontoxic, biodegradable and derived from a renewable source polymers have been studied. Among the polymers in this category, poly(lactic acid) (PLA) has been identified as a good candidate to partially substitute petroleum-derived polymer such as polypropylene, polystyrene, or polyethylene(terephthalate) in some uses[1]. Typical PLA medical applications include suture materials, drug delivery systems and oral and orthopaedic internal fixation devices. PLA is one of the commercially available biodegradable polymers from the family of aliphatic polyesters, which are produced from lactic acid, a monomer that can be synthesized from many renewable resourses such as corn and sugar beets. Althouth it has good mechanical properties such as high strength, thermoplasticity, transparency, and fabricability, its applications are limited because of its brittleness and nonflexibility[2].

In order to try to improve mechanical and thermal polymer properties and/or add new functional properties, there is a growing interest in the development of polymeric nanocomposites filled with layered compounds.

There is a wide variety of both synthetic and natural crystalline fillers that are able, under specific conditions, to intercalate a polymer[3].

Layered compounds of natural or synthetic origin are a special class of compounds in which the crystals are built by stacking of two dimensional units (the layers), which are bound to each other through weak forces. Depending on the genesis, conditions and chemical composition, the empty crystallographic sites between the layers can be occupied by anhydrous or solvated cationic, anionic, or neutral species, producing the intercalation compounds[4].

One simple classification of the layered compounds can be made according to the charge of the layers: i) Compounds with neutral layers such as graphite, layered single hydroxides like Mg(OH)2 and Al(OH)3; different transition metal chalcogenides like 2H-MoS2, 1T-TiS2, V2O5, MoO3; ii) Compounds with positively charged layers, which are compensed by the intercalation of hydrated anions as can be observed in layered double hydroxides and layered hydroxide salts; iii) Compounds, which have negatively charged layers as transition metal dichalcogenides after chemical or electrochemical reduction or some 2:1 clay minerals[4].

Cationic exchanger clays from 2:1 groups, like montmorillonite, have been extensively investigated as polymeric filler. Polymer nanocomposites filled with anionic exchangers having Brucite-like structure, layered double hydroxides (LDH), have attracted considerable technological and scientific interest.

In the Brucite structure, when part of divalent cations is isomorphically replaced by trivalent cations, positive

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Polímeros, 26(2), 106-114, 2016 107

charges are generated. In order to compensate this charge excess and stabilize the structure, exchangeable hydrated anions are inserted in the interlayer spacing, generating the LDHs. The LDHs can be represented by the general formula: [M2+

1-xM3+

x(OH)2]x+(An-)x/n.mH2O, where M2+ and

M3+ are the divalent and trivalent cations (a large number of cations combinations can be employed), An- is the interlayer anion with n- charge.

Few groups investigated the properties of the nanocomposites of PLLA filled with LDHs[1,2,5-12], but in these researches the produced LDHs were composed by only two pairs of divalent-trivalent cations: Mg/Al[1-2,5-7,9-12] and Zn/Al[8,11]. Besides, nanocomposites of PLLA filled with LDH intercalated by the anions evaluated in the present study (salicylate and sebacate) were not found.

The objective of the present study was to investigate the in situ bulk polymerization of L-Lactide in the presence of LDHs intercalated by the anions salicylate or sebacate. PLLA nanocomposites were prepared by in situ intercalative ring opening polymerization of L-lactide. LDHs containing different compositions (Zn/Al and Zn/Cr) and M+2/M+3 molar ratios (2:1) were synthesized by co-precipitation method. Mechanical, thermal and optical properties of PLLA and PLLA nanocomposites were evaluated. Although some mechanical and thermal properties of PLLA nanocomposites reinforced by LDHs have already been investigated[5,8], melt blending or solution casting routes were used to prepare the nanocomposites, besides flexural properties were not evaluated. As far as the authors know, studies considering nanocomposites of PLLA filled with LHDs containing Zn/Cr as divalent/trivalent cations were not found.

2. Materials and Methods

2.1 Materials

(3 S)-cis-3,6-dimethyl-1,4-dioxane-2,5-dione (lactide) (Sigma Aldrich, 98%) and tin (II) 2-ethylhexanoate (stannous octoate) (Sigma Aldrich, 95%) were used to produce PLLA and PLLA nanocomposites. The chemical used to prepare the LDHs like sodium hydroxide, zinc chloride, aluminum chloride, salicylic acid, sebacic acid (Ecibra), chromium chloride hexahydrate (Vetec) were of analytical grade and used as received.

2.2 Methods2.2.1 Preparation of layered double hydroxides

Layered double hydroxides were produced by co-precipitation method. This method consists in a controlled precipitation by addition of an alkaline solution[13]. In order to avoid the presence of carbonate anions between the LDH layers, this synthesis was performed under nitrogen atmosphere. After weighing the required amounts of selected chlorides, intercalating anion precursors and sodium hydroxide, they were dissolved in deionized and decarbonated water. The intercalating anion precursor solutions were placed in the reactor. The mixed salts and sodium hydroxide solutions were simultaneously and dropwise added to the reactor to maintain the pH near 10. The final mixture was left to react at a pH close to 8 for 12 hours, under dynamic flow of nitrogen. The mixture was centrifuged at 4000 rpm for

12 minutes, the supernatant was discarded and the precipitate was washed with deionized water, these processes were repeated for five times. The solids were dried in an oven for seven days at 45°C.

Zn/Al 2:1 and Zn/Cr 2:1 LDH were synthesized, where the intercalated anions were salicylate and sebacate. These anionic species were intercalated into the interlayer spacing of the layered compounds in order to try to reduce electrostatic interactions and also the hydrophilicity of the layers. Thus, the compatibility between the layered compounds and the polymeric matrix can be enhanced, making easier a possible exfoliation.

2.2.2 Preparation of PLLA nanocomposites

PLLA nanocomposites were produced by in situ intercalative ring opening polymerization using stannous octoate as catalyst. Prior to polymerization, the respective LDH (1 or 2 wt%) was mixed with L-lactide and catalyst, using a mechanical stirrer Fisatom model 713, and this mixture was vigorously stirred at 1000 rpm for 1 hour, at room temperature. Subsequently, the beaker containing the mixture was charged in a home-made borosilicate glass reactor, whose temperature was maintained by means of a Fisatom heating mantle model 67. The polymerization was carried out at 120°C for 7 hours under dynamic flow of nitrogen. The L-lactide:catalyst molar ratio was fixed at 500:1 for all the polymerization reactions. The samples were ground using an IKA A11 grinder and passed through a 100 mesh sieve.

2.3 Characterizations

The X-ray diffraction measurements (XRD) were performed using a Shimadzu-XRD 7000 diffractometer, using CuKα radiation (λ=1.5406Å), at a rate of 2°/min, operating at 40KV and 30 mA, over 2θ range of 1.5-70°.

The transmittance of the nanocomposites was examined by a spectrophotometer Varian/Cary 5G equipment in the wavelength range of 200 to 800 nm.

For the previous analyses, it was not necessary any sample preparation.

The FTIR spectra were obtained in a Spectrum One Perkin Elmer equipment, in a wavenumber range of 400 to 4000 cm-1, with a resolution of 4 cm-1, using KBr disc method.

Thermogravimetric analysis (TGA) was performed on a Universal V2.3C TA Instrument, where the samples were heated from 30 to 700°C, with a heating rate of 20°C/min on oxidant atmosphere (oxygen rate of 100 ml/min).

Differential scanning calorimetric analysis (DSC) were carried out using the DSC 2910 from TA Instruments. The samples were run from 25 to 200°C, with a heating rate of 10°C/min, held at 200°C for 5min. Then, they were cooled to -50°C, with a cooling rate of 30°C/min, held at -50°C for 5 min and, after that, the samples were heated again to 200°C with a heating rate of 10°C/min. The cooling and heating runs were performed in a nitrogen atmosphere.

In order to carry out the mechanical tests, the samples were injection molded using a Haake Minijet injection molding system from Thermo Fisher Scientific. The injection molding

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Polímeros, 26(2), 106-114, 2016108

conditions adopted were: cylinder temperature of 200°C, injection pressure of 200 bar, injection time of 10 seconds, mold temperature of 43°C, holding pressure of 150 bar and a holding time of 15 seconds. The specimens had 64 mm in length, 12.7 mm in width and 3.2 mm in thickness.

DMA analyses were performed using a dynamic mechanical analyzer from Netzsch, type DMA 242, using the holder for three-point free bending mode. The tests were done at a frequency of 1 HZ, at a force of 1N and at oscillation amplitude of 15 micrometers. The analyses were carried out at a heating rate of 5°C/min and in the temperature range of -20 to 200°C. The flexural properties were measured using a MTS testing machine model 810, considering a three-point loading system. The testing speed was 5 mm/min and span length was 30 mm.

3. Results and Discussions

3.1 Layered double hydroxides

Figures 1 and 2 show the XRD patterns, in the range of 2 theta from 1.5 to 30°, for the layered compounds produced. LDHs showed series of basal reflections and each of them represents a crystallographic plane present on the crystal structure of these layered compounds. The very broad basal reflections indicate a very small number of stacked layers which represent submicrometric crystal sizes. Zn/Cr salicylate has even smaller crystals, whose XRD pattern resembles an amorphous phase.

Through XRD technique, it was possible, by using Bragg’s law, to determine the basal distance of the studied layered materials, and the results are shown in Figures 1 and 2. X-ray diffraction patterns were interpreted with respect to the position of the basal reflection (003), which corresponds to the distance between two adjacent metal hydroxide layers in the LDH 3R polymorph lattice[14]. All the LDHs synthesized exhibited basal spacing higher than the value reported in literature to hidrotalcite (7.6 Å)[14] (natural LDH intercalated by carbonate anions), which is consistent with the intercalation of the organic anions. Figure 1 shows XRD patterns for the layered compounds intercalated by salicylate anions. The LDHs Zn/Al and Zn/Cr exhibited basal spacing of 18.24 Å and 17.8 Å, respectively.

As Brucite layer thickness is around 4.8 Â[4], these LDHs showed interlayer spacing of 13.44 Å and 13 Å, respectively. Some authors[15] reported basal spacing of 17.4 Å and 15.4 Å for Zn/Al salicylate and Zn/Cr salicylate, respectively. These differences in basal spacing can be related to different anions arrangement in the interlayer spacing. Relative to the inorganic layers, the anions can be packed in monolayer and/or bilayer, they also can be arranged perpendicularly, horizontally, or tilted at some angle and even be associated with different amounts of water. When a bilayer packing is formed, the anions can exist in an interdigitated tilted arrangement. The double layers of anions can also be tilted with an angle that can cause little or no overlap.

Figure 2 shows X-ray diffraction patterns of the LDHs(sebacate). Zn/Al and Zn/Cr LDHs exhibited basal spacing of 18.65Å and 17.38 Å, respectively. In the XRD pattern of Zn/Al, the presence of zinc sebacate is also evidenced by a peak at 12.68Å (indicated by an asterisk).

The Zn/Cr LDH exhibited broad basal peaks attributed to small crystals and basal spacing value consistent with literature data[16]. In the case of both LDHs, the basal distance is consistent with a single layer arrangement of the sebacate anions between the LDH layers.

Figure 3 shows the FTIR spectra of the LDHs studied. All the LDHs exhibited a very broad band in the range 3000-3700 cm-1. This band is LDHs characteristic and it corresponds to stretching vibration of the hydroxyl groups. LDHs(sebacate) FTIR spectra revealed bands close to 2920 cm-1 and 2850 cm-1, which can represent symmetric and asymmetric stretching vibrations of the C-H groups, respectively. LDHs exhibited bands in the range 1590-1520 cm-1 and peaks close to 1400 cm-1, which can correspond to asymmetric and symmetric stretching vibrations of the carboxylate ion, respectively. LDHs(salicylate) revealed bands close to 1370 cm-1, which can indicate the presence of carbonate anions in the interlayer region. The band close to 1270 cm-1 and 1150 cm-1 can be associated to C-O stretching vibrations.

Figure 4 shows thermogravimetric curves (TGA) for the oxidative thermal decomposition of the intercalating anion precursors and LDHs produced. When the thermal stability of the intercalating agents was compared, it was noticed that sebacic acid exhibited the highest decomposition

Figure 1. X-ray diffraction patterns: (a) salicylic acid, (b) Zn/Al salicylate, (c) Zn/Cr salicylate.

Figure 2. X-ray diffraction patterns: (a) sebacic acid, (b) Zn/Al sebacate, (c) Zn/Cr sebacate.

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Polímeros, 26(2), 106-114, 2016 109

temperature. Sebacic acid showed a final decomposition temperature close to 450°C. Figure 4 revealed that the acids evaluated in this study were submitted to almost complete degradation and showed less than 1% of char residues. Compared to salicylic acid, LDHs(salicylate) exhibited higher thermal stability and degradation residues. Salicylic acid sample suffered thermal decomposition until approximately 350°C with a char residue of 0.68%, however, at 215°C, this sample had only 1.5% of its initial weight. The thermal decomposition of Zn/Al salicylate and Zn/Cr salicylate occurred until 453°C (burning residue of 66%) and 492°C (burning residue of 54%), respectively. Sebacic acid exhibited thermal decomposition until close to 473°C, with char residue of 0.9%. Zn/Al sebacate and Zn/Cr sebacate suffered thermal degradation until 498°C (burning residue of 46%) and 319° (burning residue of 56%), respectively.

3.2 Polymeric nanocomposites

Figures 5 and 6 show X-ray diffraction patterns for PLLA/(LDH salicylate) and PLLA/(LDH sebacate), respectively. PLLA revealed a strong crystalline peak close to 16.7° (2 theta). This peak corresponds to the (200) and/or (110) plane of typical orthorhombic crystal[6]. X-ray

diffraction patterns for the PLLA/(LDH sebacate) and PLLA/(LDH salicylate) exhibited a behavior very similar to that obtained for pure PLLA. This fact can indicate that, despite the LDHs presence, a similar XRD pattern is obtained. The characteristics X-ray diffractions peaks of the LDHs(sebacate) and LDHs(salicylate) were absent, which can be explained in two ways: either there was a exfoliation of the layered compounds in the polymeric matrix or the filler content is too small to be detected, the last being the most plausible one.

Figure 7 illustrates FTIR spectra of PLLA/LDHs(salicylate) nanocomposites. As FTIR spectra of PLLA/LDHs(sebacate) nanocomposites were very similar to the ones obtained for PLLA/LDHs salicylate and to avoid repetition, they will not be shown. Bands close to 3010 cm–1 and 2950 cm–1 can be attributed to C-H stretching vibrations. The band that appears around 1760 cm–1 can be associated to C=O stretching vibration. Bands close to 1465 cm–1 and 1388 cm–1 can correspond to C-H asymmetric and symmetric bending vibrations, respectively. Bands close to 1182 cm–1 and 1090 cm–1 can be related to C-O stretching vibrations. Two bands related to the crystalline and amorphous phases

Figure 3. FTIR spectra of LDHs: (a) Zn/Al salicylate, (b) Zn/Cr salicylate, (c) Zn/Al sebacate, (d) Zn/Cr sebacate.

Figure 4. TGA curves to intercalating anion precursors: (A) Salicylic acid, (B) sebacic acid, and LDHs produced: (C) Zn/Al salicylate, (D) Zn/Cr salicylate, (E) Zn/Al sebacate, (F) Zn/Cr sebacate.

Figure 5. X-ray diffraction patterns of PLLA (a) and PLLA nanocomposites (b-e). (b) 1 wt% of Zn/Al salicylate, (c) 1 wt% of Zn/Cr salicylate, (d) 2 wt% of Zn/Al salicylate, (e) 2 wt% of Zn/Cr salicylate.

Figure 6. X-ray diffraction patterns of PLLA (a) and PLLA nanocomposites (b-e). (b) 1 wt% of Zn/Al sebacate, (c) 1 wt% of Zn/Cr sebacate, (d) 2 wt% of Zn/Al sebacate, (e) 2 wt% of Zn/Cr sebacate.

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Polímeros, 26(2), 106-114, 2016110

of PLA were found at 871 cm–1 and 756 cm–1. The band at 871 cm–1 can be assigned to the amorphous phase and the band at 756 cm–1 to the crystalline phase[17]. Bands close to 1220 cm–1 can be associated to C-O asymmetric stretching vibrations, characteristics of PLLA spectra.

Figures 8 and 9 show visible and ultraviolet transmission spectra of PLLA and PLLA nanocomposites. The ultraviolet and visible light transmission is one of the main factors that should be evaluated in polymers used for food packaging applications, in order to select materials that can preserve the foods quality until the moment of their consumption. Sensitive components of foods such as lipids, flavors, vitamins, and pigments may undergo degradation reactions when exposed to light[18]. The main driver is the need for extending the shelf-life and reducing the damages of the foods. This way, besides the suitable selection of the materials for the production of food packaging, it is also important to verify other strategies to increase polymers light barrier properties, like the incorporation of additives and/or light absorbers. Ultraviolet light and visible spectrum range from 100-400 nm and from 400-700 nm, respectively.

PLLA exhibited an increase in the UV light transmission at 234 nm and, at 318 nm, this polymer showed the highest light transmittance (95%). In the visible region, PLLA transmittance ranged from 73% to 83%. The transparency

of plastic sheeting is defined as the transmission of visible light on the range 540-560 nm[17]. In the present study, PLLA transmitted, in this range, 78% of the light. Figure 8 revealed that almost all the LDHs(salicylate) nanocomposites transmitted less UVB and UVA light than PLLA. The nanocomposite containing 1wt% of Zn/Al salicylate exhibited, on the range 540-560 nm, a PLLA similar light transmission (78%).This nanocomposite showed, at UVA and UVB wavelength region, transmittances that ranged from 51 to 81% and from 63 to 82%, respectively. PLLA exhibited, at UVA and UVB wavelength region, transmittances that ranged from 83 to 94% and from 76 to 93%, respectively. Compared to PLLA, the nanocomposite containing 1 wt% of Zn/Al salicylate transmitted less UVA and UVB light and exhibited a similar transparency. Figure 9 shows ultraviolet and visible transmission spectra of the PLLA/(LDHs/sebacate) nanocomposites. Nanocomposites light transmission spectra behavior was similar to PLLA, in the entire studied wavelength. Compared to PLLA, the nanocomposites containing Zn/Cr sebacate transmitted less UVA, UVB and visible light. These nanocomposites showed transmittances close to 55%, in the range 540-560 nm.

The results demonstrated that the addition of 1 wt% of Zn/Al salicylate in the PLLA composition can reduce the polymer transmission in the UVA and UVB wavelength region, while keeping a similar PLLA transparency. By means of these results it can be noticed that the correct selection of layered compounds can affect the PLLA light transmission.

Table 1 summarizes the main results of the oxidative thermal decomposition processes of PLLA and PLLA nanocomposites. It can be noticed that all the nanocomposites showed a T10 (temperature at which 10% of mass loss occurs) lower than PLLA. According to some authors[19] the thermal decomposition of PLLA occurs by random chain scission or specific chain-end scission, because the aliphatic ester structure is relatively easy to hydrolyze and break down. The impurities and/or the organic molecule, which have been used to modify the filler, first decompose, which provokes the subsequent thermal decomposition of PLLA. At 50% weight loss (T50), compared to PLLA, the nanocomposites containing 1 and 2 wt% of Zn/Al salicylate, revealed a slight increase on decomposition temperatures. These same nanocomposites exhibited, at 90% weight loss (T90), a small

Figure 7. FTIR spectra of PLLA (a) and PLLA nanocomposites (b-e). (b) 1wt% of Zn/Al salicylate, (c) 2 wt% of Zn/Al salicylate, (d) 1 wt% of Zn/Cr salicylate, (e) 2 wt% of Zn/Cr salicylate.

Figure 8. Percent transmission versus wavelength for PLLA and PLLA/LDH(salicylate) nanocomposites.

Figure 9. Percent transmission versus wavelength for PLLA and PLLA/LDH(sebacate) nanocomposites.

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Polímeros, 26(2), 106-114, 2016 111

increase on thermal stability, compared to PLLA. These results demonstrated that, at higher temperatures, some inorganic filler were more able prevent the thermal decomposition of PLLA. This fact can be attributed to shielding effect of layer of these LDHs to the evolution of produced products during the later stage of decomposition[19], improving the thermal stability.

Figure 10 and 11 show the storage modulus and loss modulus as a function of the temperature for PLLA salicylate and PLLA sebacate nanocomposites, respectively. Figure 10 displays that, with the exception of the nanocomposite containing 2wt% of Zn/Cr salycilate, all other nanocomposites exhibited storage modulus values very close to the ones found for PLLA. On the other hand, all the nanocomposites containing LDH sebacate in their composition showed lower elastic modulus than PLLA. The decrease in storage modulus may be caused by some agglomerated LDH platelets in the PLLA matrix that act as sites of stress concentration[5].

It was not possible to evaluate the dynamic mechanical properties of the nanocomposite containing 2 wt% of Zn/Al salicylate, because these samples were fractured at the beginning of the tests.

The glass transition temperature (Tg) can be obtained through the peak of the curve of loss modulus as a function of the temperature. The values of Tg for PLLA and the nancomposites containing 1wt% of Zn/Al salicylate, 1 wt% of Zn/Cr salicylate and 2 wt% of Zn/Cr salicylate, were of 61.7°C, 67.6°C, 62.3°C and 62°C, respectively. The nanocomposites containing 1wt% of Zn/Al sebacate, 2wt% of Zn/Al sebacate, 1 wt% of Zn/Cr salicylate and 2 wt% of Zn/Cr salicylate exhibited Tg values of 66.8°C, 63.1°C, 68.4°C and 68°C, respectively. The results displayed that, compared to PLLA, all the nanocomposites exhibited higher Tg values. These results can be attributed to physical and/or chemical interactions that may arise between the filler and the polymer[20], these interactions could reduce the available free volume, and thus the chains mobility.

At temperatures close to 50°C, the storage modulus and loss modulus were high and low, respectively, and the samples acted as a glassy material, and thus, stiff[21]. At temperatures above 80°C, both the modulus were low, and the samples exhibited a viscous behavior. When the storage modulus decreases with the increase of the temperature and the loss modulus exhibits a maximum, the viscoelastic behavior is reached.

Table 2 displays a summary of the main parameters (modulus of elasticity, maximum bending stress and elongation at break point) obtained from flexural test. The modulus of elasticity is a parameter that expresses the stiffness of a material, at the beginning of the flexural test. Through Table 2, it can be noticed that all the studied nanocomposites exhibited lower modulus of elasticity than PLLA and, as the reinforcement agent concentration increased, the modulus of elasticity decreased, for most of the samples.

In thermoplastic-based (intercalated or exfoliated) nanocomposites, the stress at break, which expresses the

Table 1. Temperatures for the thermal degradation of PLLA and PLLA nanocomposites.Sample T10 (°C) T50 (°C) T90 (°C)

PLLA 246 277 294PLLA/(1 wt% of Zn/Al salicylate) 245 282 302PLLA/(2 wt% of Zn/Al salicylate) 239 279 299PLLA/(1 wt% of Zn/Cr salicylate) 235 269 284PLLA/(2 wt% of Zn/Cr salicylate) 234 270 290PLLA/(1 wt% of Zn/Al sebacate) 234 272 295PLLA/(2 wt% of Zn/Al sebacate) 232 272 294PLLA/(1 wt% of Zn/Cr sebacate) 236 270 289PLLA/(2 wt% of Zn/Cr sebacate) 234 272 291

Figure 11. Storage modulus and loss modulus as a function of the temperature for PLLA and PLLA/(LDH sebacate) nanocomposites.

Figure 10. Storage modulus and loss modulus as a function of the temperature for PLLA and PLLA/(LDH salicylate) nanocomposites.

Nogueira, T., Gonçalves, N., Botan, R., Wypych, F., & Lona, L.

Polímeros, 26(2), 106-114, 2016112

ultimate strength that the material can bear before break, may vary strongly depending on the nature of the interactions between the matrix and the filler[3].

All the nanocomposites exhibited lower stress at break values than PLLA. These results can indicate that weak interfacial interactions between the polymer and the fillers took place. If a strong interface between the reinforcement agent and the polymer occurs, when a load is applied, it is efficiently transferred from surface to the entire polymeric matrix, by means of this interface, enhancing the mechanical properties of the nanocomposite. On the other hand, when a weak interface is present, the load distribution become less efficient, and the mechanical properties exhibit no improvement or they are negatively affected.

PLLA and the nanocomposites exhibited stress versus strain curves qualitatively similar. In these cases, the post-peak curves go down very rapidly almost in a straight line without increasing in strains. This indicates that the specimen breaks into two pieces when the maximum stress is reached[22].

Through Table 2, it can be seen that all the samples showed lower elongation at break values than PLLA. It can also be noticed that, for most of the nanocomposites, an increase in the filler concentrations caused a decrease in the strain at break values.

Table 3 summarizes the main results obtained by differential scanning calorimetry (DSC) (glass transition temperature, melting temperature, cold crystallization temperature, cold

crystallization enthalpy, melting enthalpy, percent crystallinity, melt crystallization temperature, melt crystallization enthalpy). In the present study, firstly, the samples were heated from 25°C to 200°C, held at this temperature for five minutes, in order to extinguish their thermal history. As pure PLLA and all the nanocomposites displayed both thermal transitions (glass transition temperature and melting temperature) they can be classified as semicrystalline materials.

The degree of crystallinity of the samples was calculated according to the following Equation 1:

0 *100(1 )

mc

m

HXH

∆=

−ϕ ∆ (1)

Where: Xc is the percent crystallinity, ΔHm is the melt enthalpy of the studied material, ϕ is the fraction of filler in the sample, ΔHm is the melt enthalpy of 100% crystalline PLLA that was taken as 93 J/g[12].

Through Table 3, it can be noticed that all the samples exhibited lower glass transition temperatures than PLLA. According to some authors[8] this decrease could have been due to the plasticizing effect of the anions, which were used in the intercalation of the fillers.

In this study, it was possible to obtain the crystallization temperature from molten samples (Tmc) (during the cooling scans) and from cold samples (Tcc) (during the heating scans). During the second heating scan, the nanocomposites containing 1wt% of Zn/Al sebacate and 1 wt% of Zn/Cr

Table 3. Glass transition temperature (Tg), cold crystallization temperature (Tcc), cold crystallization enthalpy(ΔHcc), melting temperature (Tm), melting enthalpy (ΔHm) percent crystallinity (Xc), melt crystallization temperature (Tmc), melt crystallization enthalpy (ΔHmc) for PLLA and the PLLA nanocomposites.

Tg(°C)(+/-0.1°C)2° heating

scan

Tcc(°C)(+/-0.1°C)2° heating

scan

ΔHcc (J/g) (+/-1%)

2° heating scan

Tm(°C)(+/-0.1°C)2° heating

scan

ΔHm (J/g)(+/-1%)

2° heating scan

Xc (%)Tmc(°C)

(+/-0.1°C)Cooling scan

ΔHmc

(+/-1%)(J/g)

Cooling scanPLLA 54.3 91.7 28.6 170.2 81.2 87.3 94 27.21 wt% of Zn/Al salicylate 49 94.8 42.4 164.7 79.8 86.7 87 14.462 wt% of Zn/Al salicylate 48.6 91.4 40.8 163 73.1 80.2 81.3 10.411 wt% of Zn/Cr salicylate 49.6 94.5 48 165.2 76.6 83.1 83 72 wt% of Zn/Cr salicylate 47.9 88.1 53.5 162.2 74.1 81.3 80 3.81 wt% of Zn/Al sebacate 39 86.1 4.1 158.1 83.2 90.3 100.9 59.172 wt% of Zn/Al sebacate 46.9 92.6 47.5 165.3 71 77.9 82.5 5.61 wt% of Zn/Cr sebacate 52.6 97.9 51.6 168.8 83.6 90.8 87.8 5.82 wt% of Zn/Cr sebacate 50.6 92.2 50.6 165.4 82.4 90.5 85.9 5.3

Table 2. Summary of the some parameters obtained from flexure test.Modulus of elasticity (MPa) Stress at break (MPa) Elongation at break (%)

PLLA 3815 +/- 147.9 20.6 +/- 4.16 0.54 +/- 0.09PLLA/(1 wt% Zn/Al salicylate) 3588 +/- 20.9 11.7 +/- 1.97 0.33 +/- 0.05PLLA/(2 wt% Zn/Al salicylate) 3434 +/- 411.1 5.8 +/- 0.25 0.17 +/- 0.015PLLA/(1 wt% Zn/Cr salicylate) 2965 +/- 187.8 8.5 +/- 0.73 0.28 +/- 0.039PLLA/(2 wt% Zn/Cr salicylate) 3158 +/- 191.1 6.7 +/- 1.39 0.21 +/- 0.033PLLA/(1 wt% Zn/Al sebacate) 3513 +/- 429.7 13.6 +/- 2.28 0.39 +/- 0.02PLLA/(2 wt% Zn/Al sebacate) 3266 +/- 123.3 7.1 +/- 0.47 0.22 +/- 0.007PLLA/(1 wt% Zn/Cr sebacate) 3420 +/- 269.5 10.7 +/- 2.78 0.3 +/- 0.062PLLA/(2 wt% Zn/Cr sebacate) 3018 +/- 202.6 6.1 +/- 0.96 0.2 +/- 0.019

Layered double hydroxides as fillers in poly(l-lactide) nanocomposites, obtained by in situ bulk polymerization

Polímeros, 26(2), 106-114, 2016 113

sebacate showed, respectively, the lowest and the highest crystallization temperatures. For other hand, during the cooling scan, the nanocomposites containing 2wt% of Zn/Cr salycilate and 1 wt% of Zn/Al sebacate exhibited, respectively, the lowest and the highest crystallization temperatures. According to some authors[8] usually, a lower Tcc indicates faster crystallization, whereas a lower Tmc indicates slower crystallization. Therefore, the decrease in Tcc coupled with the increase in Tmc could have been an indicator of a crystallization-promoting effect of the nanofiller. The nanocomposite containing 1 wt% of Zn/Al sebacate exhibited the above-mentioned conditions. This is may be to an increase in the interaction between these filler particles and the matrix, improving the nucleating activity of the filler and promoting faster crystallization[23].

Table 3 shows that all the nanocomposites exhibited lower melting temperatures than PLLA. This finding might be due to variations in the molecular weight of PLA since it is known that the PLA melting point is influenced by polymer chain length[12]. Some authors[1,10,12] reported the decrease in molecular weights of PLLA nanocomposites containing LDHs. This reduction in molecular weight could arise due to release of tightly bound water in the LDH interlayer spacing which would cause hydrolytic degradation of PLA chains. However the role of metal centers in the LDH or hydroxyl groups on the LDH surfaces as sites for catalysis of degradation cannot be disconsidered[12].

After analyzing the data of melting enthalpy, it was noticed that, compared to PLLA (ΔHm=81.2 J/g), only the nanocomposites containing 1 wt% of Zn/Al sebacate (ΔHm=83.2 J/g), 1wt% of Zn/Cr sebacate (ΔHm=83.6 J/g) and 2 wt% of Zn/Cr sebacate (ΔHm=82.4 J/g) displayed a little increase of this parameter.

The nanocomposite containing 1 wt% of Zn/Cr sebacate showed the highest percent crystallinity value. For other hand, the lowest value of this parameter was found for the nanocomposite containing 2 wt% of Zn/Al sebacate. These observations can indicate that some fillers exhibited, at certain concentrations, higher nucleation tendencies, giving rise to more uniform crystal sizes[1].

4. Conclusions

An experimental study about in situ bulk polymerization of L-Lactide filled by layered double hydroxides was carried out. It was noticed that the nanocomposite containing 1 wt% of Zn/Al salicylate exhibited the lowest transmittance in the UVA and UVB range, while keeping a transparence similar to PLLA. TGA results showed that it was possible to obtain a little increase in the decomposition temperature of the PLLA (8°C) by means of the addition of 1wt% of Zn/Al salicylate. Results of flexural test showed that all the nanocomposites presented lower stress at break and elongation at break values than PLLA. DSC results demonstrated that, compared to PLLA, all the nanocomposites displayed decreases on glass transition temperatures as well as on melting temperatures.

5. Acknowledgements

The authors would like to acknowledge the financial support from Fapesp (Fundação de amparo a pesquisa do Estado de São Paulo) [Foundation for Research Support of the State of São Paulo] and Capes (Coordenação de aperfeiçoamento de pessoal de nível superior) [Coordination for the Improvement of Higher Education Personel].

6. References

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2. Mahboobeh, E., Yunus, W., Hussein, Z., Ahmad, M., & Ibrahim, N. (2010). Flexibility improvement of poly(lactic acid) by stearate-modified layered double hydroxide. Journal of Applied Polymer Science, 118(2), 1077-1083. http://dx.doi.org/10.1002/app.32461.

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4. Wypych, F., Arizaga, G. G. C., & Satyanarayana, K. G. (2008). Synthetic layered materials/ polymer nanocomposites. In S. Thomas & G. Zaikov (Eds.). Polymer nanocomposite research advances (pp. 95-143). New York: Nova Science Publishers.

5. Chiang, M., & Wu, T. (2012). Preparation and characterization of melt processed poly(L-lactide)/layered double hydroxide nanocomposites. Composites. Part B, Engineering, 43(7), 2789-2794. http://dx.doi.org/10.1016/j.compositesb.2012.04.040.

6. Chiang, M., Chu, M., & Wu, T. (2011). Effect of layered double hydroxides on the thermal degradation behavior of biodegradable poly(L-lactide) nanocomposites. Polymer Degradation & Stability, 96(1), 60-66. http://dx.doi.org/10.1016/j.polymdegradstab.2010.11.002.

7. Chiang, M., & Wu, T. (2010). Synthesis and characterization of biodegradable poly(L-lactide)/layered double hydroxide nanocomposites. Composites Science and Technology, 70(1), 110-115. http://dx.doi.org/10.1016/j.compscitech.2009.09.012.

8. Dagnon, K., Ambadapadi, S., Shaito, A., Ogbomo, S., DeLeon, V., Golden, T., Rahimi, M., Nguyen, K., Braterman, P., & D’Souza, N. (2009). Poly(L-lactic acid) nanocomposites with layered double hydroxides functionalized with ibuprofen. Journal of Applied Polymer Science, 113(3), 1905-1915. http://dx.doi.org/10.1002/app.30159.

9. Ha, J., & Xanthos, M. (2010). Novel modifiers for layered double hydroxides and their effects on the properties of polylactic acid composites. Applied Clay Science, 47(3-4), 303-310. http://dx.doi.org/10.1016/j.clay.2009.11.033.

10. Gerds, N., Katiyar, V., Koch, C., Hansen, H., Plackett, D., Larsen, E., & Risbo, J. (2012). Degradation of L-polylactide during melt processing with layered double hydroxides. Polymer Degradation & Stability, 97(10), 2002-2009. http://dx.doi.org/10.1016/j.polymdegradstab.2012.04.014.

11. Wang, D., Leuteritz, A., Wang, Y., Wagenknecht, U., & Heinrich, G. (2010). Preparation and burning behaviors of flame retarding biodegradable poly(lactic acid) nanocomposite based on zinc aluminum layered double hydroxide. Polymer Degradation & Stability, 95(12), 2474-2480. http://dx.doi.org/10.1016/j.polymdegradstab.2010.08.007.

12. Katiyar, V., Gerds, N., Koch, C., Risbo, J., Hansen, H., & Plackett, D. (2010). Poly L-lactide-layered double hydroxide nanocomposites via in situ polymerization of L-lactide. Polymer Degradation & Stability, 95(12), 2563-2573. http://dx.doi.org/10.1016/j.polymdegradstab.2010.07.031.

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13. Arizaga, G., Satyanarayana, K., & Wypych, F. (2007). Layered hydroxide salts: synthesis, properties and potential applications. Solid State Ionics, 178(15-18), 1143-1162. http://dx.doi.org/10.1016/j.ssi.2007.04.016.

14. Costa, F., Leuteritz, A., Wagenknecht, U., Jehnichen, D., Haubler, L., & Heinrich, G. (2008). Intercalation of Mg-Al layered double hydroxide by anionic surfactants: preparation and characterization. Applied Clay Science, 38(3-4), 153-164. http://dx.doi.org/10.1016/j.clay.2007.03.006.

15. Meyn, M., Beneke, K., & Lagaly, G. (1990). Anion-exchange reactions of layered double hydroxides. Inorganic Chemistry, 29(26), 5201-5207. http://dx.doi.org/10.1021/ic00351a013.

16. Pang, X., Ma, X., Li, D., & How, W. (2013). Synthesis and characterization of 10-hydroxycamptothecin-sebacate-layered double hydroxide nanocomposites. Solid State Sciences, 16, 71-75. http://dx.doi.org/10.1016/j.solidstatesciences.2012.10.008.

17. Auras, R., Harte, B., & Selke, S. (2004). An overview of polylactides as packaging materials. Macromolecular Bioscience, 4(9), 835-864. http://dx.doi.org/10.1002/mabi.200400043. PMid:15468294.

18. Auras, R., Lim, L., Selke, S. E. M., & Tsuji, H. (2010). Poly(lactic acid) synthesis, structures, properties, processing, and applications. New Jersey: John Wiley & Sons. http://dx.doi.org/10.1002/9780470649848.

19. Chen, H., Chen, J., Shao, L., Yang, J., Huang, T., Zhang, N., & Wang, Y. (2013). Comparative study of poly(L-lactide) nanocomposites with organic montmorillonite and carbon

nanotubes. Journal of Polymer Science. Part B, Polymer Physics, 51(3), 183-196. http://dx.doi.org/10.1002/polb.23182.

20. Nogueira, T., Botan, R., Neto, J., Wypych, F., & Lona, L. (2013). Effect of layered double hydroxide, on the mechanical, termal, and fire properties of poly(methyl methacrylate) nanocomposites. Advances in Polymer Technology, 32(S1), E660-E674. http://dx.doi.org/10.1002/adv21309.

21. Cassu, S., & Felisberti, M. (2005). Comportamento dinâmico-mecânico e relaxações em polímeros e blendas poliméricas. Quimica Nova, 28(2), 255-263. http://dx.doi.org/10.1590/S0100-40422005000200017.

22. Yan, L., Chouw, N., & Yuan, X. (2012). Improving the mechanical properties of natural fibre fabric reinforced epoxy composites by alkali treatment. Journal of Reinforced Plastics and Composites, 31(6), 425-437. http://dx.doi.org/10.1177/0731684412439494.

23. Oiu, W., Mai, K., & Zeng, H. (2000). Effect of silane-grafted polypropylene on the mechanical properties and crystallization behavior of talc/polypropylene composites. Journal of Applied Polymer Science, 77(13), 2974-2977. http://dx.doi.org/10.1002/1097-4628(20000923)77:13<2974::AID-APP22>3.0.CO;2-R.

Received: July 07, 2015 Revised: Sept. 09, 2015

Accepted: Sept. 21, 2015

http://dx.doi.org/10.1590/0104-1428.2124

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Curaua fiber reinforced high-density polyethylene composites: effect of impact modifier and fiber loading

Jaqueline Albano de Morais1, Renan Gadioli1 and Marco-Aurelio De Paoli1*

1Instituto de Química, Universidade Estadual de Campinas - UNICAMP, Campinas, SP, Brazil*mdepaoli@iqm.unicamp.br

Sbstract

Short fibers are used in thermoplastic composites to increase their tensile and flexural resistance; however, it often decreases impact resistance. Composites with short vegetal fibers are not an exception to this behavior. The purpose of this work is to produce a vegetal fiber reinforced composite with improved tensile and impact resistance in relation to the polymer matrix. We used poly(ethylene-co-vinyl acetate), EVA, to recover the impact resistance of high density polyethylene, HDPE, reinforced with Curauá fibers, CF. Blends and composites were processed in a corotating twin screw extruder. The pure polymers, blends and composites were characterized by differential scanning calorimetry, thermogravimetry, infrared spectroscopy, scanning electron microscopy, tensile mechanical properties and Izod impact resistance. EVA used as impact modifier in the HDPE matrix exhibited a co-continuous phase and in the composites the fibers were homogeneously dispersed. The best combination of mechanical properties, tensile, flexural and impact, were obtained for the formulations of composites with 20 wt. % of CF and 20 to 40 wt. % of EVA. The composite prepared with 20 wt. % EVA and containing 30 wt. % of CF showed impact resistance comparable to pure HDPE and improved tensile and flexural mechanical properties.

Keywords: fibers, impact behavior, mechanical testing, extrusion.

1. Introduction

Vegetal fibers are largely replacing other reinforcing agents for thermoplastic and thermoset composites, because they are produced from renewable resources, have high toughness, have low density and are biodegradable[1]. Their use produces composites with increased tensile and flexural mechanical properties in comparison to the matrix polymer. Additionally, thermoplastic composites with vegetal fibers cause less wear to the processing equipment in comparison to glass fibers[2,3].

The vegetal fibers used in this work, Curauá fibers, CF, extracted from the leaves of an Amazonian plant of the bromeliad family (Ananas erectfolius L. B. Smith) are produced on a large scale. Like other lignocellulosic materials, are composed of hemicellulose, cellulose and lignin. We and other authors, previously characterized these fibers for their mechanical properties and chemical composition[4,5]. Their specific tensile mechanical properties are similar to those of glass fibers, making them a good candidate for their substitution[2]. The diameter of the pristine CF is in the range of 30 to 100 μm and each fiber is composed of a bundle of microfibers, which are fibrillated upon extrusion, as demonstrated in a previous work[6]. The extent of fibrillation and fiber breakage is strongly dependent on the processing conditions[7].

Cellulose acetate[8], polypropylene[9], post-consumer polypropylene[10], polyamide-6[11] and high-density polyethylene[6] were reinforced with CF in our group by processing in a corotating twin-screw extruder. In all cases, the tensile and flexural properties were improved and 20 wt. % of CF was the best content in the composites to achieve this improvement. In most cases, the use of a

coupling agent provided a better fiber to matrix interaction. Like in other polymer/fiber composites, tensile and flexural mechanical properties improvement is a consequence of stress transference to the fiber, provided by a good fiber to matrix adhesion. However, this improvement occurs in parallel to a decrease in the impact resistance, as generally observed for short fiber reinforced thermoplastic polymers[12].

High-density polyethylene, HDPE, has a high impact resistance, however, when reinforced with vegetal fibers this resistance decreases by ca. 30%[13]. The decrease in impact resistance was also observed for composites of high-density biopolyethylene with Curauá fibers, in different proportions and processed by two different methods[14]. To attain the impact resistance of the neat biopolyethylene the same authors used liquid hydroxylated polybutadiene, as compatibilizer and impact modifier, in the composite with Curauá fibers. Similar impact resistance reduction and use of impact modifiers was reported for short jute fiber reinforced polypropylene composites[15].

The most common strategy to improve the impact resistance, i.e. toughness, of a thermoplastic material is to make a blend with an elastomer. This is routinely used on an industrial scale, like when polycarbonate becomes tougher by blending with high impact polystyrene (Noryl™ is the commercial name of this blend). To obtain the tougher material it is necessary to make an immiscible and compatible blend, controlling the concentration of the modifier, particles size, distance between the particles and degree of adhesion between the polymeric phases[16]. In this case, the rubbery phase concentrates or absorbs tension, changing the tension

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Polímeros, 26(2), 115-122, 2016116

states of the matrix. The tension absorbed by the modifier causes plastic deformation, which absorbs the impact.

Previous works report on the rheology, morphology and mechanical properties of polyethylene blended with poly(ethylene-co-vinyl acetate) in different compositions and using various comonomer contents[17-19]. These works focused on the miscibility of the blends but did not report on the effect of EVA content on the impact resistance of the blend.

As discussed above, vegetal fiber composites of HDPE have improved tensile and flexural mechanical properties in comparison to pure HDPE; however, the drawback is the reduction of the impact resistance. Thus, in the area of thermoplastic biocomposites the existing gap is a material that combines high tensile and flexural resistance with good impact resistance. Making the composite with a blend of HDPE with an immiscible softer material can recuperate this property. Thus, the aim of the present work is to produce a composite material associating the high tensile and flexural strength of the composite of HDPE with CF and the good impact resistance of pure HDPE. This is pursued by using EVA as impact modifier for the composite of HDPE with CF.

2. Materials and Methods

2.1 Material used

CF supplied by Embrapa-PA (Belém, Brazil) were fractioned in a three-knife rotary mill (Rone, model NFA 1533), with knife/counter-knife gap of 0.5 mm, washed two times with tap water, dried in open air for 24 h and further dried in a conventional oven at 100 °C for 8 h. Before processing, the fibers were disaggregated in a mechanical homogenizer (MH Equipamentos) for 2 min, to facilitate the feeding into the side-feeder of the extruder.

HDPE, JV060U, was supplied by Braskem (Triunfo, Brazil) with melt flow index of 6.1 - 8.0 g 10 min-1 at 190 °C and 2.16 kg and ρ = 0.957 g cm-3. EVA, HM2528, was from Braskem (Triunfo, Brazil) with 28 wt. % of vinyl acetate and melt flow index of 25 g 10 min-1 at 190 °C and 2.16 kg.

2.2 Processing

Blends and composites were processed in a corotating intermeshing twin-screw extruder (Coperion Werner & Pfleiderer, Germany, ZSK 26 Mc, L/D = 44) with degassing and a side-feeder. Feeding was controlled with two Brabender DDW-MD2-DSR28-10 gravimetric dosimeters. For composites processing we used a previously published screw design[20], while, for the blends, part of the kneading elements were replaced by transport elements to reduce shear. The composition of the blends and composites is shown in weight percent, wt. %, in Table 1.

Blends were prepared by pre-mixing 20, 30 or 40 wt. % of EVA into HDPE and extrusion in the following conditions: temperature profile of 170, 175, 180, 185, 190, 190, 195, 200, 200, 195 and 200 ºC, from feed to die, and screw rotation of 300 rpm. The average torque and mass temperature were 30% and 208 °C, respectively.

Composites of all the HDPE/EVA blends were processed by extrusion with 20 or 30 wt. % of fibers using a temperature profile of 120, 120,125, 125, 130, 130, 135, 135, 140, 135 and 130 °C, from feed to die, and screw rotation of 300 rpm. Fibers were fed using the side-feeder, at 265 rpm, and vacuum degassing was employed in all processes. The average torque and mass temperature were 47% and 151 °C, respectively.

Test samples with dimensions according to ASTM D-638 and ASTM D-256 were prepared by injection-molding (Arburg, All-rounder M250) using previously dried (100 ºC for 1 h) pellets of the blends and composites. For the blends the conditions were: temperature profile of 180, 185, 190, 200 and 205 °C, 1.4 x 108 Pa of injection pressure, 1.0 x 108 Pa of hold pressure, injection velocity of 15 cm3 s-1, mold temperature of 20 °C and 20 s of cooling time. For the composites with 20 wt. % of fibers the conditions were: temperature profile of 150,160, 170, 175 and 165 °C, 1.2 x 108 Pa of injection pressure, 8.0 x 107 Pa of hold pressure, injection velocity of 15 cm3 s-1, mold temperature of 20 °C and 10 s of cooling time. For composites with 30 wt. % of fibers the conditions were: temperature profile of 170, 180, 190, 200 and 210 °C, 1.5 x 108 Pa of injection pressure, 1.3 and 1.4 x 107 Pa of hold pressure, injection velocity of 3 cm3 s-1, mold temperature of 20 °C and 10 s of cooling time.

2.3 Characterization of blends and composites

Differential scanning calorimetry, DSC, was used to determine the melting temperature, Tm, and glass transition temperature, Tg, using a DSC-Q100 equipment (TA Instruments) with a cooling and heating rate of 10 °C cm-1 in the temperature range of -50 to 200 °C, under an argon atmosphere (50 mL min-1).

Thermogravimetric analyses, TGA, was used to determine the mass loss parameters using a TA2900 (TA Instruments) in the temperature range from 20 to 600 °C, with a heating rate of 10 °C cm-1 under synthetic air flow (100 mL min-1).

Scanning electron microscopy, SEM, was used to monitor the morphology of the cryogenic fracture surface of the samples. For this, the injection molded test samples were maintained for 30 min in liquid N2 before fracturing. The surface of the fractures were coated with carbon and gold by sputtering using a Balzers MD BalTec 020 equipment. SEM analyses were done at an accelerating voltage of 5 kV using a Jeol JSM–6360LV equipment.

Table 1. Composition of blends and composites in wt. %. HDPE = high density polyethylene, EVA = poly(ethylene-co-vinyl acetate) and CF = Curauá fiber.

Acronym HDPE EVA CFHDPE 100 - -HDPE /20EVA 80 20 -HDPE/30EVA 70 30 -HDPE/40EVA 60 40 -HDPE/20EVA20CF 64 16 20HDPE/30EVA20CF 56 24 20HDPE/40EVA20CF 48 32 20HDPE/20EVA30CF 56 14 30

Curaua fiber reinforced high-density polyethylene composites: effect of impact modifier and fiber loading

Polímeros, 26(2), 115-122, 2016 117

Fourier transform infrared, FTIR, spectra in ATR mode were obtained using a Smiths IlluminatIR II-Micro-Infrared Probe for Optical Microscopy, range 4000 - 400 cm-1, 21 scans min-1 and resolution of 4 cm-1. The spectra of pure HDPE and blends with EVA were obtained by ATR using samples in the form of films, produced by pressing in a hydraulic press (Marconi MA-098-A) at 160 °C and 0.02 MPa for 2 min., For the composites, the samples were milled to observe the contribution of the fibers to the spectra measured as KBr pellets in the transmittance mode.

For mechanical properties determinations, the samples were conditioned for 48 h in an acclimatized room at 25 (± 5) °C and 50 (± 5) % relative humidity. The tensile properties were determined according to ASTM D-638 using an EMIC DL2000 equipment with a 5 kN load cell. For the impact resistance we used the ASTM D-256 standard and an EMIC AlC1 equipment with a 2.7 J hammer.

3. Results and Discussions

3.1 HDPE/EVA Blend

3.1.1 Differential Scanning Calorimetry

From the DSC curves of the blends in the second heating scan, Figure 1a, we determined the melting temperatures, Tm, as the peak of the endothermic transition. The Tm of the HDPE phase in the blends did not change with increased contents of EVA, in relation to pure HDPE. The first heating is routinely used in DSC measurements to erase the thermal history of the sample. These curves also show that there was no shift of the glass transition temperature (Tg) of the EVA phases in the blends, as highlighted in Figure 1b for the HDPE/40EVA blend. The Tg of the HDPE phase was not detected in the blends because it is below the temperature range used (ca. - 95 °C). These DSC results demonstrate that, according to the Tg criteria, EVA and HDPE form immiscible phases in these blends. By determining the area under the melting peak, it is concluded that the presence of EVA does not change the degree of crystallinity of the HDPE phase.

3.1.2 Thermogravimetry

In Figure 2, we compare the thermogravimetric curves of HDPE, EVA and the blend HDPE/20EVA, measured under synthetic air flow. Other blend compositions showed similar behaviors. The onset mass loss in these curves (ca. 250 oC) corresponds to the oxidative thermal degradation onset of pure HDPE, the polyethylene block of EVA and the HDPE phase of the blend. The inflections observed in the pure HDPE curve are a consequence of the formulation of the grade used. The TGA curve of EVA reaches a plateau at ca. 70% mass, with a subsequent sharp drop to 15% at 440 oC, corresponding to the vinyl acetate block oxidative thermal degradation with the formation of acetic acid and conjugated double bonds[21]. Residue formation above 450 oC is similar for all three samples. The blend curve suggests that there is no interaction between the thermo oxidation reactions of the blend components.

3.1.3 Infrared spectroscopy

Figure 3 illustrates the reflectance FTIR spectra of HDPE, EVA and the blend HDPE/20EVA (all blend compositions were measured) measured directly at the surface of the injection-molded specimens. The infrared spectrum provides information about structural aspects of polymers, such as chemical composition, conformation and structural configuration. The position and relative intensities of the absorption bands may give information concerning the interaction between the blend components. We observe that the reflectance infrared spectra of the blends correspond exactly to a superimposition of the spectra of its components with no band shifts. This is another indication for the lack of chemical interaction between the blend components, corroborating the conclusion that it is an immiscible blend.

3.1.4 Scanning Electron Microscopy

Figure 4 shows the SEM micrographs of the cryofractured surface of injection molded samples of the HDPE/EVA blends with 30 and 40 wt. % of EVA. Figure 4a shows the micrograph of the crude fracture and in Figure 4b the EVA phase was extracted with acetone to highlight the morphology.

Figure 1. a) DSC curves (2nd heating) of HDPE blends with 20, 30 and 40 wt. % of EVA and (b) detail of the part of the curve used for Tg determination in a) for HDPE/40EVA. Curves in a) were vertically shifted for better comparison.

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As shown in the micrographs, the HDPE/EVA blends exhibit a co-continuous morphology, with EVA as dispersed phase and HDPE as continuous phase. This phase separation

confirms the immiscibility of the blend components. SEM micrographs of the other blends were also obtained and show the same features. The extraction of EVA with acetone resulted in “empty” spaces homogenously distributed in the HDPE matrix, Figure 4b, confirming that EVA is uniformly distributed in the HDPE phase. The transmittance FTIR spectrum of a film cast from the acetone extract corresponds to the spectrum of EVA, confirming its extraction.

3.1.5 Mechanical properties

We measured the tensile properties of the blends to determine the effect of the EVA concentration on the tenacity. Figure 5 shows the average stress-strain curves of pure HDPE and the blends. Figure 6 compares the tensile strength and modulus for the blends and pure HDPE (values shown in Table 3). The approximately linear decrease of tensile strength and modulus with EVA concentration in the blends also reflects their immiscibility.

The tenacity in Table 2 was calculated as the area under the stress-strain curves. Comparing to pure HDPE, there is a large increase in tenacity in the HDPE/20EVA blend. This value slightly decreases for HDPE/30EVA and significantly drops when the concentration of EVA is increased in the HDPE/40EVA blend. This indicates that 30 wt. % of EVA is the upper limit of concentration to increase tenacity of HDPE with an EVA copolymer containing 28% of vinyl acetate.

3.2 HDPE/EVA blends reinforced with Curauá fibers

3.2.1 SEM

The SEM micrographs of the cryogenic fracture of the injection molded samples of the composites of the HDPE blend with 20, 30 and 40 wt. % EVA and 20 wt. % CF are depicted in Figure 7. In Figure 7a we observe that the incorporation of the fibers in the blend did not change its morphology, with a co-continuous phase similar to the pure blend, shown in Figure 4a. Figure 7b illustrates the homogeneous distribution of the fibers in the blend and their adhesion to the matrix, evidenced by low occurrence of fiber

Figure 2. Comparison of the thermogravimetric curves of HDPE, EVA and the blend HDPE/20EVA, measured under synthetic air flow at 10 oC min-1.

Figure 3. ATR-FTIR spectra for EVA, HDPE, and the blend of HDPE/20EVA. Spectra vertically shifted for comparison.

Figure 4. SEM micrographs of the fracture surface of injection molded blends: (a) HDPE/30 EVA with scale bar of 1 µm and (b) HDPE/40EVA after extraction with acetone with scale bar of 10 µm.

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pullout during fracture. The good fiber-matrix adhesion shown in Figs 7a and highlighted by a circle, confirms the good interaction between the blend and the reinforcing agent. Higher magnification (not shown) reveals that the fibers are partially covered by the polymer. The micrograph of the cryogenically fractured composite extracted with acetone is shown in Figure 7d with higher magnification. Some voids observed between the fibers and the matrix in Figure 7d can be an indication that the EVA phase is responsible for the good fiber/matrix adhesion in the blend. This argument is reinforced by the absence of holes left by EVA extraction in Figure 7d, when compared to Figure 4b.

The diameter of the pristine Curauá fibers is ca. 100 µm and it is reduced by fiber fibrillation, occuring as a consequence of the high shear of the twin-screw extruder, as reported in previous works[6]. The presence of microfibrils in the SEM micrographs, with diameters of the order of 5 µm or less, observed in Figures 7c and 7d, confirm that the CF fibers were fibrillated during processing. This diameter reduction reflects in a high aspect ratio of the fibers, improving the reinforcement effect[22].

3.2.2 Infrared spectrophotometry

Figure 8 shows the transmittance infrared spectra of the composites of the HDPE/EVA blend with 20, 30 and 40 wt. % of EVA and 20 wt. % of Curauá fibers. These spectra were measured after dispersing the milled composites in KBr. This was done because the reflectance spectra of the surface would show no evidence for the presence of the fibers, which tend to migrate to the bulk of the material during injection molding[12]. Bands labeled 1 to 8 correspond to HDPE and EVA in the blend, particularly band 6 at 1746 cm-1, corresponding to the C=O stretching of the acetate group of the EVA phase and band 1 at 727 cm-1, corresponding to the CH2 bending of HDPE and EVA. The presence of the fibers in the composites is evidenced by the bands labeled as a, b and c. The most evident is the broad band centered at 3440 cm-1, band c, corresponding to the O-H stretching of the cellulose and lignin hydroxyl groups, present in the fibers[4].

3.2.3 Thermogravimetry

Figure 9 shows the TGA curve for the composite of the HDPE blend with 20 wt. % of EVA and 20 wt. % of CF, HDPE/20EVA20CF. The curve shows two main degradation processes starting at ca. 250 and 400 °C. The first is assigned

Table 3. Tensile and impact resistance results for pure HDPE, HDPE/EVA blends and blends reinforced with Curauá fiber. (* from ref.[6]).Sample Tensile

Strength(MPa)

Strain atBreak(%)

Young’sModulus

(%)

Impact Resistance(MPa)

Pure HDPE* 18.0 ± 0.2 - 543 ± 55 91 ± 3HDPE/20CF* 28.7 ± 0.4 3.7 ± 0.2 1375 ± 147 57 ± 3HDPE/20EVA 13.3 ± 0.1 376 ± 100 368 ± 20 #

HDPE/30EVA 11.7 ± 0.2 386 ± 71 297 ± 61 #

HDPE/40EVA 10.3 ± 0.2 338 ± 23 231 ± 18 #

HDPE/20EVA/20CF 23 ± 1 7 ± 1 866 ± 38 128 ± 3HDPE/30EVA/20CF 20 ± 1 7 ± 2 546 ± 87 132 ± 11HDPE/40EVA/20CF 16.9 ± 0.2 20 ± 3 585 ± 70 140 ± 11HDPE/20EVA/30CF 23.0 ± 0.3 5.6 ± 0.2 1273 ± 26 88 ± 3# - samples did not break.

Figure 5. Stress-strain curves for HDPE and the blends containing 20, 30 and 40 wt. % of EVA.

Figure 6. Comparison of tensile strength and Young modulus for the HDPE/EVA blends, as a function of EVA content.

Table 2. Tenacity, calculated as the integrated area under the stress-strain curves for HDPE and the blends.

Formulations Area (MPa mm)HDPE 229

HDPE/20EVA 1964HDPE/30EVA 1818HDPE/40EVA 1384

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to thermal decomposition of the fiber, by comparing to other previously prepared composites[6], and the second to the blend, as compared to Figure 2. Similar to previous

results reported in the literature[23], good fiber to matrix interaction favors the stability of the composite, shifting the matrix onset degradation temperature to a higher value.

Figure 7. SEM micrographs of the cryogenic fracture surface of injection molded composites: (a) HDPE/20EVA20CF showing fiber-matrix interaction and co-continuous EVA phase, (b) HDPE/30EVA20CF showing the homogeneous distribution of the fibers, (c) HDPE/40EVA20CF highlighting fiber diameter and (d) HDPE/40EVA20CF composite after extraction with acetone with higher magnification.

Figure 8. Transmittance infrared spectra of the composites of HDPE/EVA blends, with different EVA contents, and 20 wt. % of CF, measured as KBr pellets. Spectra were vertically shifted for comparison.

Figure 9. TGA curve for the composite HDPE/20EVA20CF, under synthetic air flow.

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Polímeros, 26(2), 115-122, 2016 121

The absence of residue above 520 oC indicates absence of mineral contaminants in the fibers. The first derivative curve (not shown) indicates no interaction between the degradation products of the fibers and the blend. The TGA curves of the remaining composites show similar features.

3.2.4 Tensile properties

Figure 10 and Table 3 present the variation of the mechanical properties of the composites with 20 and 30 wt. % of CF and toughened with different concentrations of EVA, in comparison to pure HDPE, to its composite with 20 wt. % of CF[6] and blends with different EVA contents. As observed previously, there is a strong increase of tensile strength and modulus by dispersing 20 wt. % of fibers in HDPE. When making the composites with the blends and 20 wt. % of fibers, these values decrease with the increased concentration of EVA in the blend and are recovered by increasing fiber content to 30 wt. %. Thus, the composite with 50 wt. % of HDPE, 20 wt. % of EVA and 30 wt. % of CF (HDPE/20EVA/30CF) shows the same tensile strength and a higher modulus in comparison to HDPE/20EVA/20CF, which contains 60 wt. % of HDPE. Thus, increasing the fiber content compensates for the effect of EVA on the tenacity of the matrix.

3.2.5 Impact resistance

Figure 10b and Table 3 show the variation of the impact resistance of the composites with 20 wt. % of CF and 20, 30 or 40 wt. % of EVA and for the composite with 30 wt. % of CF and 20 wt. % of EVA (HDPE/20EVA/30CF). It is clear that all composites with the HDPE/EVA blend have a higher impact resistance in comparison to pure HDPE. This is due to the EVA phase of the blend, which is responsible for the impact force dissipation. Additionally, increasing the fiber content in the 20EVA composite to 30 wt. % (HDPE/20EVA/30CF) yields a material with impact resistance similar to pure HDPE. Note that in this composite the HDPE content is only 50 wt. %.

4. Conclusions

HDPE and EVA form an immiscible blend with a high impact resistance in comparison to pure HDPE, because the dispersed EVA phase absorbs the impact energy. By using this blend to make composites with Curauá fibers, it is possible to obtain a material with good tensile resistance and, simultaneously, an impact resistance comparable to that of pure HDPE. This demonstrates that the mechanical properties of vegetal fibers based thermoplastic composites can be tailored for high impact resistant material applications. It is also important to note that the composite with the best mechanical properties uses 70% of petrochemical based raw materials and 30 wt. % of a renewable resource, with clear environmental benefits.

5. Acknowledgements

The authors thank EMBRAPA-PA for the fibers, Braskem for the polymer samples and FAPESP for financial support (Proc. 2010/17804-7). JAM thank SAE-Unicamp for a scholarship. We also thank Profs. Maria do Carmo Gonçalves and Maria Isabel Felisberti for suggestions.

6. References

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4. Tomczak, F., Satyanarayana, K. G., & Sydenstricker, T. H. D. (2007). Studies on lignocellulosic fibers of Brazil: Part III – Morphology and properties of Brazilian curauá fibers. Composites. Part A, Applied Science and Manufacturing, 38(10), 2227-2236. http://dx.doi.org/10.1016/j.compositesa.2007.06.005.

Figure 10. Comparison of the mechanical properties of pure HDPE, HDPE/20CF composite and composites of blends with different EVA and CF contents in wt. %: a) tensile strength and Young’s modulus and b) impact resistance.

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5. Spinacé, M. A. S., Lambert, C. S., Fermoselli, K. K. G., & De Paoli, M. A. (2009). Characterization of lignocellulosic Curaua fibres. Carbohydrate Polymers, 77(1), 47-53. http://dx.doi.org/10.1016/j.carbpol.2008.12.005.

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9. Mano, B., Araujo, J. R., Waldman, W. R., Spinacé, M. A. S., & De Paoli, M. A. (2013). Mechanical properties, morphology and thermal degradation of a biocomposite of polypropylene and Curauá fibers: coupling agent effect. Polímeros: Ciência e Tecnologia, 23, 161-168. http://dx.doi.org/10.4322/S0104-14282013005000025.

10. Spinace, M. A. S., Fermoselli, K. K. G., & De Paoli, M. A. (2009). Recycled polypropylene reinforced with Curauá fibers by extrusion. Journal of Applied Polymer Science, 112(6), 3686-3694. http://dx.doi.org/10.1002/app.29683.

11. Santos, P. A., Spinacé, M. A. S., Fermoselli, K. K. G., & De Paoli, M. A. (2009). Efeito da forma de processamento e do tratamento da fibra de curauá nas propriedades de compósitos com poliamida-6. Polímeros: Ciência e Tecnologia, 19(1), 31-39. http://dx.doi.org/10.1590/S0104-14282009000100010.

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13. Spinace, M. A. S., Janeiro, L. G., Bernardino, F. C., Grossi, T. A., & De Paoli, M. A. (2011). Polyolefins reinforced with short vegetal fibers: Sisal vs. Curauá. Polímeros: Ciência e Tecnologia, 21, 168-174. http://dx.doi.org/10.1590/S0104-14282011005000036.

14. Castro, D. O., Ruvolo-Filho, A., & Frollini, E. (2012). Materials prepared from biopolyethylene and curaua fibers: composites

from biomass. Polymer Testing, 31(7), 880-888. http://dx.doi.org/10.1016/j.polymertesting.2012.05.011.

15. Rana, A. K., Mandal, A., & Bandyopadhyay, S. (2003). Bandyopadhyay S.: Short jute fiber reinforced polypropylene composites: effect of compatibilizer, impact modifier and fiber loading. Composites Science and Technology, 63(6), 801-806. http://dx.doi.org/10.1016/S0266-3538(02)00267-1.

16. Rabello, M., & De Paoli, M.-A. (2013). Aditivação de termoplásticos. São Paulo: Artliber.

17. Faker, M., Razavi Aghjeh, M. K., Ghaffari, M., & Seyyedi, S. A. (2008). Rheology, morphology and mechanical properties of polyethylene/ethylene vinyl acetate copolymer (PE/EVA) blends. European Polymer Journal, 44(6), 1834-1842. http://dx.doi.org/10.1016/j.eurpolymj.2008.04.002.

18. Khonakdar, H. A., Wagenknecht, U., Jafari, S. H., Hässler, R., & Eslami, H. (2004). Dynamic mechanical properties and morphology of polyethylene/ethylene vinyl acetate copolymer blends. Advances in Polymer Technology, 23(4), 307-315. http://dx.doi.org/10.1002/adv.20019.

19. Khonakdar, H. A., Jafari, S. H., Yavari, A., Asadinezhad, A., & Wagenknecht, U. (2005). Rheology, Morphology and estimation of interfacial tension of LDPE/EVA and HDPE/EVA blends. Polymer Bulletin, 54(1-2), 75-84. http://dx.doi.org/10.1007/s00289-005-0365-6.

20. Araujo, J. R., Waldman, W. R., & De Paoli, M. A. (2008). Thermal properties of high density polyethylene composites with natural fibres: Coupling agent effect. Polymer Degradation & Stability, 93(10), 1770-1775. http://dx.doi.org/10.1016/j.polymdegradstab.2008.07.021.

21. McNeill, I. C. (1997). Thermal degradation mechanisms of some addition polymers and copolymers. Journal of Analytical and Applied Pyrolysis, 40-41, 21-41. http://dx.doi.org/10.1016/S0165-2370(97)00006-5.

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23. Albano, C., Gonzalez, J., Ichazo, M., & Kaiser, D. (1999). Thermal stability of blends of polyolefins and sisal fiber. Polymer Degradation & Stability, 66(2), 179-190. http://dx.doi.org/10.1016/S0141-3910(99)00064-6.

Received: Feb. 20, 2015 Revised: Sept. 14, 2015

Accepted: Jan. 27, 2016

http://dx.doi.org/10.1590/0104-1428.2195

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Effect of compatibilization in situ on PA/SEBS blendsAnna Paula Azevedo de Carvalho1 and Alex da Silva Sirqueira1*

1Laboratório de Tecnologia de Materiais, Universidade Estadual da Zona Oeste – UEZO, Rio de Janeiro, RJ, Brasil

*alexsirqueira@uezo.rj.gov.br

Sbstract

This paper presents the mechanical, thermal and rheological properties of polymer blends of polyamide 6 (PA) and styrene-butylene-styrene (SEBS) using SEBS-g-MA as a coupling agent. The increase in the interfacial interaction of PA/SEBS blends with the addition of SEBS-g-MA was observed to enhance the mechanical properties studied (excellent elongation at break revealed by a 500% increase). Higher hardness values and a higher degree of crystallinity were obtained for compatibilized blends. In the presence of the SEBS-MA compatibilizer, the heat sweep thermograms obtained by DSC showed only one melting peak, which confirms the effect of the compatibilizer. The best rheological properties were observed for the ternary blends. The lowest concentration of SEBS-g-MA promoted the toughening of PA, and 7.5 phr of compatibilizer increased the modulus E’ of the ternary blends compared to that of the binary blends. DMTA analysis allowed for the system to be characterized as partially miscible.

Keywords: polyamide, SEBS, compatibilization, rheology, DMTA.

1. Introduction

Polymer blends enjoy widespread application in various industries for the development of new materials capable of combining the intrinsic properties of each component and can be obtained by the simple method of physical blending. This class of materials represents an economically viable alternative for various applications because of its low production cost[1]. Examples of materials in this class are those with a high impact strength consisting of heterophasic systems such as high-impact polystyrene and mixtures of polypropylene and elastomers[2].

Although the materials obtained from polymer blends afford economic advantages, the resulting mixtures of these systems exhibit weaker properties. The synergism of properties is rarely achieved due to various factors such as structural differences, differences in polarity, and the rate of crosslinking among different materials. These effects produce an immiscible blend without the final properties desired. The properties of an immiscible polymer blend are dependent on factors such as the morphology of the system, composition, the rheological and thermal properties of the components, and processing conditions. To improve the performance of heterophasic systems, it is common to use fillers, reinforcing agents and compatibilizers. The addition of small amounts of a third component in the immiscible blend can change the interfacial energy and the dispersion between phases, acting as a compatibilizer[3-4].

Coupling agents are the most widely used non-reactive agents, as they allow for the reaction control required for reactive compatibilizers. Non-reactive agents are usually block or graft and can lead to a reduction in size of the dispersed phase and modifying the mechanical properties of the material. These agents must be compatible with each polymer in a mixture. Selecting the appropriate coupling agent is a great challenge because the choice depends on factors such as molecular weight, chemical structure and interaction

between each component in a homopolymer/copolymer mixture[5].

One of the methods used to obtain compatible polymer blends involves the modification of polymers with maleic groups that can be used as coupling agents. The thermoplastic elastomer composed of styrene and butadiene functionalized with maleic anhydride (SEBS-g-MA) has been widely used for compatibilizing polyamides with polypropylene[6].

The objective of the present study was to develop a thermoplastic elastomer based on polyamide 6 (PA) and styrene-butylene-styrene (SEBS) to meet the performance requirements of the automotive industry. However, due to the immiscibility between the pair of materials, SEBS-g-MA was used as the coupling agent in the system. It is known that different types of polyamides are amino terminal groups capable of reacting easily with terminal anhydride groups. Figure 1 shows a hypothetical model proposed for the reaction between the amino end groups of PA and the anhydride functional group of SEBS that occurs in the in situ compatibilization mechanism to combine the pair.

2. Materials and Methods

2.1 Materials

The polyamide-6 (PA) used in this work was a commercial product from Radici, Italy. The melt flow index, or MFI, (at 230 °C and 2.16 kg load) and density were 35 g/10 min and 1.13 gr/cm3, respectively. Styrene-ethylene/butylene-styrene (SEBS) triblock copolymer containing 30 wt% styrene and styrene-ethylene/butylene-styrene triblock copolymer grafted with 1.84 wt% of maleic anhydride (SEBS-g-MA) were supplied by Kraton Polymers as commercial-grade Kraton G 1650 and Kraton FG1901, respectively. SEBS-g-MA has been reported to contain a styrene to ethylene/butylene ratio

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in the triblock copolymer of 30/70 by weight and an MFI value of 22 g/10 min (at 230 °C and 5 kg load).

2.2 Preparation of PA/SEBS blends

Prior to the mixing process, PA, SEBS and SEBS-g-MA were dried at 80 °C for 24 hours in a vacuum oven. PA/SEBS blends were prepared by a mechanical mixer in a Haake Rheomix OS internal mixer. Polyamide was introduced into the mixer at a set temperature of 260 °C, in a CAM rotor geometry, a rotor speed of 60 RPM, and a fill factor of 0.7; after 3 minutes, the SEBS and SEBS-g-MA were mixed for more 5 minutes under the same test conditions. Thus, a total of 8 minutes was required to obtain PA/SEBS blends in a melt state. The sample compositions are reported in Table 1.

2.3 Characterizations2.3.1 Mechanical properties

Prior to testing, all specimens were dried in an oven at 80 °C for 10 hours. Tensile tests were performed using a Universal Machine Test (EMIC model 422- Brazil) at room temperature according to ASTM D638; sample type IV was used at a cross-head speed of 100 mm/min. The Shore D hardness was recorded over 1 s after a presser foot contacted each sample, according to ASTM D2240 (Bareiss Hardness - Germany).

2.3.2 Swelling

The specimens to be tested in swelling were immersed in a bottle containing engine oil (Petrobras 15/50 W, Brasil) at 80 ºC for 70 hours, according to ASTM D316-95. Thereafter, the specimens were removed from the oil, quickly dipped in toluene, and blotted lightly with filter paper to eliminate excess oil on the specimen surfaces. The degree of swelling was calculated by Equation 1:

( ) b a

b

W WSwelling % 100W−

= × (1)

where Wb is the weight before swelling and Wa is the weight after swelling.

2.3.3 Thermal analysis

The thermal behaviours of the melting temperature (TM) and enthalpy of fusion (ΔHf) were obtained by differential scanning calorimetry (DSC) using a Maia F3 instrument. The analyses were performed under inert conditions with a N2 flow rate of 50 ml/min and a heating rate of 20 °C/min. All sample masses were approximately 5 mg. Melting thermograms were obtained after the second heating run. The degree of crystallinity (XC) of the blends was calculated from the ratio of the enthalpy of fusion of the PA/SEBS mixture and the enthalpy of fusion of 100% crystalline polyamide 6 (ΔHPA/SEBS/ΔHPA), where ΔHPA = 190,8 J.g–1[7].

2.3.4 Dynamic mechanical analysis (DMA)

The dynamic mechanical properties of the PA/SEBS blends were measured using a dynamic mechanical analyser from TA Instruments. Properties including tan delta and the storage modulus (E´) of the PA/SEBS blends were examined in the bending mode of deformation at a heating rate of 10 °C/min and a frequency of 1Hz from –130 to 130 °C.

2.3.5 Thermogravimetric analysis (TGA)

Thermogravimetric analysis was performed on a TGA instrument from TA Instruments under a nitrogen atmosphere and at a heating rate of 10 °C/min. A small amount (approximately 9 mg) of specimen was used for analysis. The weight loss of the specimen was measured as a function of temperature.

2.3.6 Rheological properties

Rheological analyses were performed on a Haake rheometer - MARS model, from Thermo Scientific Instruments using a parallel plate geometry with a plate diameter of 25 mm and a gap of 1 mm. The complex viscosity (η *), the shear storage modulus (G’) and the shear loss modulus (G”) of the binary and ternary blends were measured at 260 °C over the frequency range of 0.01 to 15 Hz with a strain of 1%. It was ensured that all experiments were performed in the linear viscoelastic regime. The complex strain sweeps were performed at a frequency of 1 Hz to determine the region of linear behaviour of the viscoelastic material. The strain rate used ranged from 0.1 to 600%.

3. Results and Discussions

3.1 Torque rheometry

Torque rheometry has frequently been used to monitor chemical reactions during reactive melt mixing. The torque-time behaviour of the PA/SEBS/SEBS-g-MA blend is shown in Figure 2. The compatibilizer (SEBS-g-MA) was added to PA/SEBS after 4 minutes of processing in an internal mixer, when the PA was duly melted and SEBS plasticized. When 5 phr of SEBS-g-MA was added to the PA/SEBS mixture, an increase in torque was observed. This effect can be attributed to the reaction between MA and the amine group in PA. According to Roeder et al.[8], Jiang et al.[9] and Bassani et al.[10], the anhydride groups of SEBS-g-MA react with the terminal amine groups of PA 6, forming an imide group and resulting in a copolymer in situ at the interface. Figure 1 illustrates a hypothetical reaction for the formation of the PA-SEBS pair. The reaction between the anhydride groups of SEBS-g-MA and terminal amine groups of PA involves the formation of water as a

Figure 1. Hypothetical reaction for the formation of PA-SEBS pair.

Table 1. Formulation of mixtures PA/SEBS.PA

(Phr)SEBS (Phr)

SEBS-g-MA (Phr)

50 50 -50 45 5.050 42.5 7.550 35 15

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Polímeros, 26(2), 123-128, 2016 125

byproduct, which can lead to the degradation of the PA chains by hydrolysis. However, Figure 2 shows that the end torque of the PA/SEBS/SEBS-g-MA blend remained constant, which indicates that no degradation occurred. Regarding the other torque curves of PA/SEBS, it should be noted that the blend containing 7.5 phr of compatibilizer showed the highest torque, indicating that a high content of SEBS-g-MA promotes a highly effective reaction with PA. The sample containing 15% SEBS-MA showed the same behaviour as that observed for the sample containing 7.5 phr of compatibilizer.

3.2 Mechanical properties

The effect of compatibilization on the mechanical properties of tensile strength was evaluated. Table 2 presents the tensile strength at break, elongation at break, MFI and hardness for the blends studied. The addition of SEBS-g-MA to the PA/SEBS blends yielded higher tensile strength and elongation at break. The tensile strength for the samples containing 5 and 7.5 phr of compatibilizer increased by approximately 100%. The melt index flow showed the highest values for the compatibilized blend, likely due to the compatibilizer’s effect in anchoring the phases, which reduced the distance between PA and SEBS. The increase in the mechanical properties studied (tensile, elongation, feasibility to flow and hardness) can be explained by an increase in molecular weight due to the compatibilization reaction, leading to a reduction in the interfacial tension between different phases and consequently the average particle size of the dispersed phase. Compared with that measured for the binary blend, a significant increase in the elongation at break for the compatibilized blends was observed, reaching a strain of 500%. This behaviour can be attributed to the improvement of the interaction between dispersed phase and matrix with the incorporation of SEBS-g-MA. In some cases, the hardness of a material defines its applicability. Table 2 presents the hardness values for the PA/SEBS blends. The compatibilized blends showed higher hardness values compared to those of the PA/SEBS blends. This behaviour was due to the incorporation of SEBS into the polyamide matrix, which was favoured by the effects of the compatibilizer. It was observed that the addition of 5 phr of SEBS-g-MA improved all mechanical properties.

The effect of the compatibilizer on the physical properties of swelling in oil is illustrated in Figure 3. The degree of swelling decreased by approximately 70% for the mixtures containing SEBS-g-MA. The addition of a compatibilizing agent to the system reduced the interfacial tension, causing an increase in the interaction between the phases. It is

possible that the compatibilizing agent has promoted the presence of a third phase or an interface that has low affinity for the oil, preventing the penetration of small molecules, which explains the decrease in the degree of swelling of the PA/SEBS/compatibilizer mixtures[11].

3.3 Thermal analysis

Figure 4 shows the thermogram obtained for the fusion of PA and the PA/SEBS PA/SEBS/SEBS-g-MA blends. The figure shows that PA has two fusion peaks. These two peaks are commonly observed in PA and are assigned to different crystalline forms, i.e., α and γ, as reported in the literature[12]. The peak with the highest TM corresponds to the crystalline structure α, and the lower peak corresponds to the γ crystalline portion. The literature[6] suggests that the calculation of both XC and TM allow for a better understanding of the results. The PA/SEBS thermograms showed two TM peaks, but the compatibilizer blend showed a tendency to form a peak, reflecting the effect of SEBS-MA in reducing the interfacial tension.

Figure 2. Curves ilustring torque rheometry as a function of time for PA/SEBS and PA/SEBS/SEBS-g-MA mixtures.

Figure 3. Properties of swelling for PA/SEBS and PA/SEBS/SEB-g-MA mixtures.

Table 2. Mechanical properties of elongation at break, tensile strength, melt index flow and hardness for the PA/SEBS and PA/SEBS/SEBS-g-MA mixtures.

PA/SEBS/SEBS-g-

MA (Phr)

Elongation at break

(%)

Tensile strength (MPa)

MFI(g/10 min)

Hardness (Shore D)

50/50/00 159 ± 49 17.9 ± 0.7 7.2 ± 5.0 52 ± 750/45/5.0 798 ± 55 32.9 ± 2.5 12.7 ± 0.9 60 ± 1

50/42.5/7.5 730 ± 40 30.5 ±0.9 13.1 ± 0.5 57 ± 150/35/15 486 ± 62 20.4 ± 1.3 12.4 ± 0.3 60 ± 2

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Polímeros, 26(2), 123-128, 2016126

Table 3 presents the thermal properties measured for each of the different compositions. The presence of SEBS-g-MA decreased TM PA; however, two melting peaks were observed for these blends, demonstrating the immiscibility of the system. For the compatibilized blends, the thermograms of heating showed only one melting peak, which confirmed the effect of the compatibilizer. The reduction in the TM values of the ternary mixtures may reflect changes in the distribution of the crystallites of PA.

The presence of SEBS-g-MA also reduced the XC of the ternary mixtures, which approached the Xc of pure PA. Marco et al.[13] performed experiments using PP/PA6/PP-MA as a compatibilizer and also reported a decrease in the XC of PA6 using a coupling agent. This behaviour can be explained by the decrease in mobility due to the grafting of PA6. This result confirms the hardness results obtained.

3.4 Dynamic mechanical thermal analysis

Dynamic mechanical thermal analysis involves applying an oscillating force on a sample and is very efficient characterizing viscoelastic polymeric systems, breaking down the module into an elastic component and a viscous component. Studying the dependence of visco-elastic properties on temperature allows for the determination of the modulus of elasticity and the damping value[14].

Figures 5 and 6 shows the thermodynamic-mechanical properties obtained by DMTA analysis for the binary and compatibilized blends. The damping properties (tan delta) are shown in Figure 5. The PA/SEBS blend (50/50) showed behaviour expected for immiscible systems,

Table 3. Thermal properties for binary and PA/SEBS/SEBS-g-MA mixtures.

PA/SEBS/SEBS-g-MA

(Phr)TM (°C) ΔH (J/g) XcPA (%)

100/00/00 233* 45.36/26.11 15*50/50/00 224 57 1650/45/5.0 218 32 17

50/42.5/7.5 221 29 1550/35/15 221 30 16

*Average between TM and XcPA corresponding to α and γ crystalline portions.

Figure 4. Melting thermograms of PA (a) and the binary and ternary mixtures (b).

Figure 5. Effect of the presence of a coupling agent on the tan delta values of binary and ternary PA/SEBS mixtures.

Figure 6. Effect of the presence of a coupling agent on the storage modulus of binary and ternary PA/SEBS mixtures.

Effect of compatibilization in situ on PA/SEBS blends

Polímeros, 26(2), 123-128, 2016 127

demonstrating two peaks, one peak corresponding to PA and another corresponding to SEBS. The PA/SEBS blends showed the Tg of PA, at a temperature close to 60 °C, and the Tg of the copolymer, close to –40 °C. The ternary blends (PA/SEBS/SEBS-g-MA) showed a third peak temperature at approximately 23 °C (located between the Tg of SEBS and Tg of PA) and an increase in the Tg of PA, which rose to 80 °C. The third peak can be explained by the presence of the SEBS-g-MA giving rise to the formation of micro-bridge stages, generating a PA-SEBS pair (see Figure 1). The strong interactions between MA along the SEBS backbone and amine groups gave rise to a decrease in damping. Despite the decrease in damping, the SEBS-MA blends also displayed a slight shift in the glass transition temperature towards lower values. The shift in Tg can be related to increased interfacial interactions and reduced slip. For the compatibilized blends, it was observed that 15 phr of SEBS-MA reduced the damping in the low-temperature region and 5 phr of SEBS-MA reduced the damping in the PA region.

The storage modulus (E’) is a very important property for evaluating the mechanical behaviour of a polymeric composite. Figure 6 shows a vitreous plateau and a reduction in the modulus E’ of the ternary mixture with 5 phr of SEBS -g-MA compared with the binary mixture. This result is due to the lower crystallinity of this sample, which favours the toughening of PA. For blends containing a higher content of compatibilizer (7.5 and 15 phr), an increase in the modulus E’ was observed. This effect can be explained by the increased interfacial interaction between the phases causing better stress distribution at the interface, which imparted greater rigidity to that sample[15]. It is interesting to note that no significant changes occurred in the blends containing a high amount of SEBS-MA (15 phr). It is likely that the system was saturated at 7.5 phr.

3.5 Rheological properties

The rheological properties of polymers are the most important for processing because they affect all processes involving the flow of material. Viscosity is the property that characterizes the rheological behaviour of polymers. The viscosity value indicates the resistance of a polymer to a given type of flow and shear that may be permanent or oscillatory. The study of the viscosity of polymer blends allows for the evaluation of, among other factors, interactions between the different phases in the blends, such as the interfacial tension[16].

Figure 7 shows the curves of the elastic modulus (G’) and viscous modulus (G”) as a function of frequency for the PA/SEBS and PA/SEBS/SEBS-g-MA blends. It should be noted that in the low-frequency region, all samples showed an elastic modulus superior to the viscous modulus. In this study, G’ increased as the amount of SEBS-MA increased, likely due to the MA and amine reactions. At high frequency, when the material response is similar to that of an elastic solid, all compatibilized blends showed the following relation G’> G”; only non-compatibilized blends showed the inverse behaviour.

This behaviour can be attributed to the increase in the molecular weight by the strong interfacial interaction caused by the compatibilization[17]; in others words, SEBS-MA acted as a reactive interfacial agent. Figure 8 shows the relationship between the viscosity and shear rate; it should be noted that all compatibilized blends showed a higher viscosity than the non-compatibilized blends in the linear viscoelastic region (LVR)

3.6 Thermogravimetry

To test the thermal stability of the non-compatibilized and compatibilized blends, thermogravimetric analysis (TGA) was performed. Figure 9 compares the thermograms of the PA/SEBS blends. The degradation of the non-compatibilized PA6/SEBS blends occurred at a lower temperature than that the compatibilized blends. The non-compatibilized PA/SEBS blends showed two degradation steps, which can be related to the PA6 and SEBS phases. However, it is interesting to note that the compatibilized blends showed a single degradation step due to the anchoring effect that occurred between the PA6 and SEBS phases, increasing the thermal stability. The high amount of SEBS-MA did not promote the thermal stability of the blends.

Figure 7. G’ and G” as a function of ω for PA/SEBS non-compatibilized and compatibilized blends.

Figure 8. Viscosity as a function of shear rate for PA/SEBS non-compatibilized and compatibilized blends.

Carvalho, A. P. A., & Sirqueira, A. S.

Polímeros, 26(2), 123-128, 2016128

4. Conclusions

The results of this study show that SEBS-g-MA acts as a compatibilizer in mixtures of polyamide and SEBS. The increase in the interfacial interaction of PA/SEBS blends with the addition of SEBS-g-MA proved to increase the mechanical properties studied (excellent elongation at break) and also increased the Tg of PA. Furthermore, higher hardness values and a higher degree of crystallinity were observed for compatibilized blends. The best rheological properties were observed for ternary mixtures. A lower concentration of SEBS-g-MA promoted the toughening of PA and 7.5 phr of compatibilizer increased the modulus E’ of the ternary blends compared to that of the binary blends. DMTA analysis allowed for the blend system to be characterized as partially miscible. Based on the results obtained for all of the properties, the best SEBS-MA concentration was 5 phr. Furthermore, this concentration improved the thermal stability of the blends in particular.

5. Acknowledgements

The authors acknowledge the support of the Laboratory of Polymer Blends and Composite Conductors (LMPCC/IMA-UFRJ) and the Laboratory of Processing and Characterization of Materials (LPCM - INT) in testing the rheological behaviour of the materials formed in this study, the CNPq, FAPERJ and Centro Universitário da Zona Oeste (UEZO) for financial support.

6. References

1. Utracki, L. A. (1989). Polymer alloys and blends: thermodynamics and rheology. New York: Hanser Publisher.

2. Pavlidou, S., & Papaspyrides, D. (2008). A review on polymer-layered silicate nanocomposites. Progress in Polymer Science, 33(12), 1119-1139. http://dx.doi.org/10.1016/j.progpolymsci.2008.07.008.

3. Paul, D. R., & Robeson, L. M. (2008). Polymer technology: nanocomposites. Polymer, 49(15), 3187-3204. http://dx.doi.org/10.1016/j.polymer.2008.04.017.

4. Maglio, G., & Palumbo, R. (1994). The role of interfacial agents in polymer Blends, Processing, Morphology and Properties. New York: Plenum Press.

5. Gomes, S. A., Barbosa, R. V., & Soares, B. G. (1992). Agentes compatibilizantes não reativos para blendas polimericas. I. síntese de poli(etileno-co-acetato de vinila-G-estireno). Polímeros: Ciência e Tecnologia, 2, 19-24. Retrieved in 21 June 2015, from http://revistapolimeros.org.br/articles/view/id/515c6c9e1ef1fae740000540

6. Fiegenbaum, F. (2007). Estudo da compatibilização das blendas PP/PA e PP/EPR (Master’s dissertation). Universidade Federal do Rio Grande do Sul, Porto Alegre.

7. Kusomoto, Ishak, M. Z. A., Chow, W. S., Takeichi, T., & Rochmadi. (2008). Influence of SEBS-g-MA on morphology, mechanical and thermal properties of PA/PP/organoclay nanocomposites. European Polymer Journal, 44, 1023-1029. http://dx.doi.org/10.1016/j.eurpolymj.2008.01.019

8. Roeder, J., Oliveira, R. V. B., Gonçalves, M. C., Soldi, V., & Pires, A. T. N. (2002). Polypropylene/polyamide-6 blends: influence of compatibilization agent on interface domains. Polymer Testing, 21(7), 815-821. http://dx.doi.org/10.1016/S0142-9418(02)00016-8.

9. Jiang, C., Filippi, S., & Magagnini, P. (2003). Reactive compatibilizer precursors for LDPE/PA6 blends. I: maleic anhydride grafted polyethylenes. Polymer, 44(8), 2144-2149. http://dx.doi.org/10.1016/S0032-3861(03)00133-2.

10. Bassani, A.; Hage, E., Jr.; Persan, L. A.; Machado, A. V., & Covas, J. A. (2005). Evolução da morfologia de fases de blendas PA6/AES em extrusora de dupla rosca e moldagem por injeção. Revista Polímeros: Ciência e Tecnologia, 15, 176-185. http://dx.doi.org/10.1590/S0104-14282005000300007

11. Almeida, M. S. M. (2006). Compatibilização reativa e vulcanização dinâmica em misturas de polipropileno e borracha nitrílica (Doctoral thesis). Universidade Federal do Rio de Janeiro, Rio de Janeiro.

12. Shi, D., Ke, Z., Yang, J., Gao, Y., Wu, J., & Yin, J. (2002). Rheology and morphology of reactively compatibilized/PA6 blends. Macromolecules, 35(21), 8005-8012. http://dx.doi.org/10.1021/ma020595d.

13. Marco, C., Ellis, G., Gómez, M. A., Fatou, J. G., Arribas, J. M., Campoy, I., & Fontecha, A. (1997). Rheological properties, crystallization, and morphology of compatibilized blends of isotactic polypropylene and polyamide. Journal of Applied Polymer Science, 65, 13-20. http://dx.doi.org/10.1002/(SICI)1097-4628(19970926)65:13<2665::AID-APP8>3.0.CO;2-9.

14. Campbell, D., Pethrick, R. A., & White, J. R. (2000). Polymer characterization: physical techniques. New York: Hanser Publisher.

15. Sanchez, A., Rosales, C., Laredo, E., Muller, A. J., & Pracella, M. (2001). In situ fibrillar reinforced PET/PA6/PA66. Polymer Science Engineer, 41, 205-208. http://dx.doi.org/10.1002/pen.10721.

16. Walters, K. (1989). An introductiuon to rheology. Amsterdam: Elsevier Press.

17. Ollshom, B., Hassander, H., & Tornell, B. (1998). Improved compatibility between polyamide and polypropylene by the use of maleic anhydride grafted SEBS. Polymer, 39, 26-33. http://dx.doi.org/10.1016/S0032-3861(97)10290-7.

Received: June 21, 2015 Revised: Sept. 07, 2015

Accepted: Oct. 08, 2015

Figure 9. TGA results obtained for PA/SEBS blends.

http://dx.doi.org/10.1590/0104-1428.2306

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Polímeros, 26(2), 129-136, 2016 129

Structure-flammability relationship study of phosphoester dimers by MLR and PLSa

Luminita Crisan1, Smaranda Iliescu1 and Simona Funar-Timofei1*

1Institute of Chemistry Timisoara, Romanian Academy, Timisoara, Romania*timofei@acad-icht.tm.edu.ro

Sbstract

Polyphosphonates and polyphosphates having good flame retardancy represent an important class of organophosphorus based polymer additives. In this analysis the flammability of 28 previously synthesized polyphosphoesters, modelled as dimmers, was explored using the multiple linear regression (MLR) and Partial Least Square (PLS) methodology. The statistical quality of the final MLR and PLS models was estimated using the following parameters: the squared correlation coefficient (r2

training = 0.917 and 0.976), the training root-mean-square errors (RMSEtr = 0.029 and 0.016) and the leave-seven-out cross-validation correlation coefficient (q2

L70 = 0.748 and 0.881), respectively. External validation was checked for a test set of seven compounds using several criteria. The MLR models had somewhat inferior fitting results. The final MLR and PLS models can be used for the estimation of limiting oxygen index (LOI) values of new polyphosphoester structures. The presence of phosphonate groups and increasing molecular branching in an isomeric series favour the dimer flammability.

Keywords: quantitative structure-property relationships, polyphosphonate, polyphosphate, limiting oxygen index, flame retardancy.

1. Introduction

An important feature of most commercial polymers is to be non-flammable or flame retardant[1]. Other polymer properties, like as: glass transition temperature, thermal decomposition temperature, etc., have been previously studied by quantitative structure-property relationships[2,3].

Flame retardant polymeric materials containing phosphorus, like poly(alkyl or aryl)phosphonates, display good flame retardancy[4].

Different polyphosphoesters with fire retardant properties were reported in the literature, being included in materials like: polycarbonates, polyamides, thermosets, etc[5]. The flammability of phosphorous polymers was investigated in order to determine structural–property relationships, too[6,7]. Two types (R and S) of chirality were found for the monomer polyphosphoesters, which were geometry optimized using the MMFF94s force field[6]. Multiple linear regression (MLR), artificial neural networks (ANNs) and support vector machines (SVMs) were applied to correlate the limiting oxygen index (LOI) values to the structural calculated descriptors. Good fitting results and predictable models were obtained using the MLR and ANN approaches, the SVM modelling providing the poorest results. It was concluded that the monomer geometry is important for flame retardancy.

Our goal was to develop robust multiple linear regression (MLR) and the partial least squares (PLS) models that select a set of variables that efficiently predict the limiting oxygen index (LOI) values and guide new information on

the flammability mechanism of polyphosphoesters[6] dimers. This parallel approach gives the opportunity to compare the quality of results supplied by the two methodologies.

2. Materials and Methods

2.1 Data set

We used a series of 28 previously synthesized polyphosphoesters[6], which were modelled in the present study as dimers. The dataset in this investigation consisted of 28 RR, RS, SR and SS phosphoester dimers for compounds 1 to 14; compounds 15 to 28 had only one chiral centre, at the P2 phosphorous atom (see Figure 1).

Experimental data for the limiting oxygen index (LOI), expressed in % (Table 1), and used as dependent variable in this study, was previous reported in references[6] and[7]. Dimer molecular structures were built using the Marvin program[8], which was used for drawing, displaying and characterizing chemical structures. Dimer conformers were pre-optimized using the 94s variant of the MMFF (Merck Molecular force field)[9] with coulomb interactions and the attractive part of the van der Waals interactions, included in the OMEGA software[10-12]. The following parameters were used for the conformer generation: a maximum of 400 conformers per compound, an energy cut-off of 10 kcal/mol relative to a global minimum identified from the search. SMILES notation was used as program input. The stereoisomers were generated using the ‘Flipper’ utility inside the Omega program. To avoid redundant conformers, any conformer having a RMSD fit outside 0.5 Å to another conformer was removed.a Dedicated to the 150th anniversary of the Romanian Academy.

Crisan, L., Iliescu, S., & Funar-Timofei, S.

Polímeros, 26(2), 129-136, 2016130

2.2 Molecular descriptor calculation

Molecular descriptors were calculated for the optimized dimer structures, using the DRAGON[13] and InstantJchem (which was used for structure database management, search and prediction)[14] software. The 1511 Dragon molecular descriptors were divided into twenty-two logical blocks, as follows: constitutional descriptors, topological descriptors (MSD-mean square distance index (Balaban), PW4-path/walk 4 - Randic shape index), walk and path counts, connectivity indices, information indices (IC5-information content index (neighborhood symmetry of 5-order)), 2D autocorrelations (Gats5e-Geary autocorrelation - lag 5/weighted by atomic Sanderson electronegativities), edge adjacency indices (EEig09d-Eigenvalue 09 from edge adj. matrix weighted by dipole moments), BCUT descriptors, topological charge indices (GGI1-topological charge index of order 1, JGI2-mean topological charge index of order2), eigenvalue based indices, Randic molecular profiles, geometrical descriptors, RDF descriptors, 3D-MoRSE descriptors (Mor15e-3D-MoRSE - signal 15/weighted by atomic Sanderson electronegativities, Mor13p-3D-MoRSE - signal 13/weighted by atomic polarizabilities Mor13m-3D-MoRSE - signal 13/weighted by atomic masses), WHIM descriptors, GETAWAY descriptors (R2m+ - R maximal autocorrelation of lag 2/weighted by atomic masses), functional group counts (nP(=O)O2R-number of phosphonates), atom-centered fragments, charge

descriptors, molecular properties, 2D binary fingerprints, and 2D frequency fingerprints. Then the molecular descriptors were verified and constant or near-constant variables were eliminated. The calculated molecular descriptors play a fundamental role in transforming the chemical information into a numerical code suitable for application in computation[15].

2.3 Training and test set selection

The series of phosphoester dimers were divided into training and test set using several approaches: the partition against medoids (PAM) algorithm[16] (“cluster” package available in R[17] based on the Euclidian distance), the decreasing response order and the random splitting. In order to use same test set in both MLR and PLS approaches, seven out of twenty eight (25%) phosphoester dimers (compounds 2, 10, 11, 15, 17, 19 and 22, see Figure 1) were chosen as test set to validate the final models. The data structures and the LOI range values (in %), comprised in the test set (0.22-0.50) and the training set (0.18-0.55), are commensurate.

2.4 Multiple Linear Regression (MLR) and Partial Least Square (PLS)

Multiple linear regression (MLR)[18] has been applied after variable selection carried out by means of a genetic algorithm included in the QSARINS v. 2.2 program[19,20] using

Figure 1. Dimer phosphoester structure. RR series: R chiral centre at P1, R chiral centre at P2; RS series: R chiral centre at P1, S chiral centre at P2; SR series: S chiral centre at P1, R chiral centre at P2; SS series: S chiral centre at P1, S chiral centre at P2; compounds 15 to 28 had only one chiral centre, at the P2 phosphorous atom.

Structure-flammability relationship study of phosphoester dimers by MLR and PLS

Polímeros, 26(2), 129-136, 2016 131

the RQK fitness function, with leave-one-out cross-validation correlation coefficient, which constrained the function to be optimized. In MLR, the number of 1549 calculated descriptors is too high compared to the number of compounds (N = 28) and an appropriate variable selection method was required. MLR calculations were carried out separately for each dataset: RR, RS, SR, SS.

In MLR calculations the structural data was normalized based on the autoscaling method, which can be described as:

mj mmj

m

X XXT

S−

= (1)

where for each variable m, XTMJ and XMJ are the j values for the m variable after and before scaling, respectively, mX is the mean and SM the standard deviation of the variable.

The PLS methodology is a generalization of the MLR one, having as main advantage the possibility to analyze the data with correlated, noise, and large number of independent variables[21]. In the PLS equation the latent variables were transformed as function of the original XIJ (i =1, 2,..., N; j=1, 2,..., K) variables, resulting following equation:

0 1 1 2 2ˆ ... ...i i i j ij k ikY b b X b X b X b X= + + + + + + (2)

where ŶI represents the calculated dependent variable, and bJ the PLS coefficients. The obtained models were optimized by a procedure of outlier detection and based on variables with significant coefficients different from zero. When the variable selection was achieved, only the significant descriptors with coefficients different from zero were preserved in the final models (for noise elimination).

Both methologies have as main goal to find out a mathematical model with minimum number of parameters and with good estimation capability.

2.5 Model validation

For the external validation of the MLR and PLS models several parameters were calculated: Q2

F1[22]

, Q2F2

[23],

Q2F3

[24] (models with values higher than 0.7 were considered acceptable), the CCCext (the concordance correlation coefficient, with satisfactory values higher than 0.85)[25], RMSEext (root-mean-square errors) and MAEext (mean absolute error)[26] and R2

pred (a higher limit than 0.5 was considered as acceptable)[27]. The comparable thresholds used in this study for different validation criteria have been rigorously previously determined[25,28]. Other statistical parameters[29] were used for the external test set: (i) squared correlation coefficient (r2

test) between the predicted and observed activities as well as squared correlation coefficient by cross-validation (q2); (ii) coefficient of determination for linear regressions with intercepts set to zero, i.e. r2

0 (predicted versus observed activities), and r′2

0 (observed versus predicted activities); (iii) slopes k and k’ of the above mentioned two regression lines. The following conditions should be satisfied for a model with acceptable predictive ability:

2 0.5q > (3)

2 0.6testr > (4)

2 20

2( ) 0.1 0.85 1.15r r and k

r−

< ≤ ≤ (5)

2 '20

2( ) 0.1 0.85 ' 1.15r r and k

r−

< ≤ ≤ (6)

2 '20 0 0.3r r− < (7)

For the internal validation of the final models other parameters were employed: r2

training (determination coefficient), q2

L70 (leave-seven-out cross-validation coefficient; values higher than 0.7 were considered as acceptable), q2

LOO (leave-one-out cross-validation coefficient), RMSEtr, MAEtr and CCCtr, calculated for the training set.

In the mean time higher r2training values must be accompanied

by q2 values as close to the r2training ones as possible[30] (to

avoid over fitting, which was, also, checked by the RMSE and MAE values).

The risk of chance correlation was, also, verified by the Y-scrambling procedure (r2

Scr and q2Scr ) and must have

lower values than the original model. For calculation of r2Scr

and q2Scr this process was repeated 999 times in case of PLS

calculations and 2000 times in the MLR ones.After the check of all validation parameters, the

applicability domain for the models is required, because robust and validated models cannot be expected to reliably predict the modelled property for any type of compounds. The applicability domain is a theoretical region in physicochemical of response and chemical structure space for which a QSAR model should make predictions with a given reliability[30]. In the applicability domain only the predictions for those compounds that fall within this domain can be considered as reliable, not extrapolations of the model. In the Williams plot the standardized residuals versus the leverages (hi) was exploited to visualize the applicability domain for our final MLR models.

3. Results and Discussions

The major objective of this paper was the estimation of limiting oxygen index (LOI) of phosphoester dimers using molecular descriptors that can be computed directly from molecular structure and guide new information on the flammability mechanism.

3.1 MLR results

The relationship between the molecular descriptors and LOI values of the dimer derivatives is illustrated by the following Equations 8-11:

RR model

2 2

0.56( 0.03) - 0.23( 0.03) -

0.19( 0.03) 09 0.20( 0.03) 20.05( 0.02) ( ) 2

0.03 0.896 44.09 0.864adj LOO

LOI MSD

EEig d R mnP O O R

SEE r F q

+

= ± ±

± + ± +± =

= = = =

(8)

RS model

2 2

0.55( 0.05) 0.18( 0.04) 4 -0.26( 0.05) 10.16( 0.06) 2 - 0.35( 0.06) 13

0.05 0.787 19.48 0.745adj LOO

LOI PWGGIJGI Mor m

SEE r F q

= ± + ±± +± ±

= = = =

(9)

Crisan, L., Iliescu, S., & Funar-Timofei, S.

Polímeros, 26(2), 129-136, 2016132

SR model

2 2

0.38( 0.02) 0.19( 0.04) 50.13( 0.05) 15 - 0.33( 0.04) 13

0.04 0.839 35.60 0.790adj LOO

LOI ICMor e Mor p

SEE r F q

= ± + ± +± ±

= = = =

(10)

SS model

2 2

0.52( 0.04)0.18( 0.04) 5 - 0.15( 0.04) 5 -0.29( 0.03) 13

0.04 0.870 45.41 0.831adj LOO

LOIIC GATS eMor p

SEE r F q

= ± +± ±±

= = = =

(11)

where SEE represents the standard error of estimates, r2adj - the

adjusted r2, F- the Fischer test, q2LOO -leave-one-out cross-validation

coefficient. Other statistical results of models 8-11 are included in Tables 2, 3, 4.

The Williams (of the standardized residuals versus the leverage) plot was used to visualize the applicability domain of the final best MLR_RR model (Figure 2). This plot confirms the absence of outliers and influential points. All compounds were located within the applicability domain and were predicted accurately.

The MLR_RR model is completely satisfactory in the fitting and has high predictive power. The LOO (leave-one-out) cross-validation highlights that the model is stable, not obtained by chance, in fact the difference between r2

training and q2

LOO is small: 5.3%. This model is internally predictive with differences between q2

LMO and q2LOO of -4.5%, and between

r2training and r2

LMO of 9.8%.The risk of chance correlation was, also, verified by

the Y-scrambling procedure. The extremely low calculated r2Scr and q2

Scr scrambling values (Table 2) indicate no chance correlation for the chosen models.

The RMSE (root-mean-square error) values for the training and validation sets are similar. The chosen MLR_RR model demonstrate a satisfactory stability in internal validation, has high fitting, internal and external predictive power.

The high values of Q2F1, Q

2F2, Q

2F3 and CCCext external

validation parameters (see section 2.5) included in Table 3

and all calculated terms of Golbraikh and Tropsha (Table 4) confirm the predictive power of all MLR models.

Better statistical results and a more stable model to simulate polymer flammability were noticed in case of the RR dataset model compared to the others.

The edge adjacency matrix encodes information about the connectivity between graph edges[15]. EEig09d (eigenvalue 09 from edge adj. matrix weighted by dipole moments) takes into account the molecular polarity, being unfavourable for dimer flame retardancy.

The mean square distance index, denoted as MSD[15], is calculated from the second-order distance distribution moment[31]. The MSD index decreases with increasing molecular branching in an isomeric series, which is favourable for dimer flammability.

GETAWAYs (Geometry, Topology, and Atom-Weights Assembly) are geometrical descriptors encoding information on the effective position of substituents and fragments in the molecular space[15]. Moreover, they are independent of molecule alignment and they, also, account to some extent for information on molecular size and shape as well as for specific atomic properties. Increased R2m+ (R maximal autocorrelation of lag 2/weighted by atomic masses) values favour the dimer flammability. Compounds containing phosphonate groups are favourable for the dimer flame retardancy.

3.2 PLS results

PLS calculations were performed with SIMCA-P+12[32] program using 21 stereoisomers as a training set and 7 stereoisomers as a test set with the taken ratio of 75% for training set and 25% for test set in whole series of compounds. The large difference between the r2

training and q2L70

values of the first calculated PLS model (lower than 0.3 is accepted) demonstrated the model over fit, and suggested the need for enhancement of the model quality. Therefore, the noise variables (with insignificant coefficient values) have been removed. Several PLS models were developed for the RR, RS, SR and SS datasets to increase their predictive power. In the final PLS_SS model compound 5 was omitted, being found as outlier, in accordance to the Hotelling’s T2 range plot[32].

The final (four-components for the RR, RS, SR datasets and two-components for the SS dataset) PLS models are satisfactory in the fitting (Table 2). The over fitting of the models was exceeded by the remarkable high and close values of r2

training and q2L70, and was, also, checked by the RMSE and

MAE (mean absolute error) parameters. In the same time similar RMSE values for the training and validation sets are observed (Tables 2 and 3).

PLS models with predictive power were obtained (see Tables 3 and 4), except the PLS_SS one, as seen from the values of Q2

F1, Q2F2, Q

2F3 and CCCext parameters. The predicted

LOI values for the RR dataset are given in Table 1.The PLS models were internally validated using, also,

999 permutations in Y-scrambling. The calculated r2Scr and q2

Scr scrambling values (Table 2) indicate no chance correlation for the chosen models.

Figure 2. Williams plot: standardized residuals of the MLR_RR model versus leverages, predicted by fitting. Training compounds are marked by white circles and test compounds by black circles.

Structure-flammability relationship study of phosphoester dimers by MLR and PLS

Polímeros, 26(2), 129-136, 2016 133

Table 1. Experimental and predicted LOI values, structural descriptors included in the final MLR_RR model.

No Exp. LOI Calc. LOI by MLR_RR

Calc. LOI by PLS_RR MSD EEig09d R2m+ nP(=O)O2R

1 0.38a 0.38 0.36 0.218 2.266 0.022 22 0.35a 0.41 0.37 0.228 2 0.023 23 0.30a 0.33 0.33 0.235 2.245 0.026 24 0.48a 0.45 0.45 0.213 2.058 0.023 25 0.55a 0.55 0.55 0.191 2.652 0.187 26 0.28a 0.26 0.28 0.235 2.569 0.043 27 0.29a 0.32 0.29 0.218 2.576 0.044 28 0.42b 0.44 0.41 0.2 2.271 0.021 29 0.44a 0.42 0.44 0.182 2.744 0.062 210 0.50a 0.45 0.54 0.212 2.467 0.117 211 0.47b 0.52 0.54 0.169 2.361 0.022 212 0.32a 0.31 0.32 0.209 2.621 0.016 213 0.40a 0.45 0.44 0.191 2.361 0.021 214 0.33a 0.28 0.31 0.219 2.619 0.019 215 0.28a 0.29 0.25 0.216 2.434 0.018 016 0.25a 0.28 0.25 0.226 2.359 0.025 017 0.22a 0.24 0.27 0.232 2.431 0.026 018 0.36a 0.31 0.37 0.211 2.454 0.025 019 0.40a 0.37 0.47 0.19 2.722 0.082 020 0.18a 0.19 0.18 0.232 2.691 0.04 021 0.20a 0.23 0.19 0.216 2.695 0.032 022 0.31a 0.35 0.32 0.198 2.434 0.017 023 0.33a 0.37 0.35 0.181 2.749 0.063 024 0.50a 0.49 0.50 0.21 2.514 0.204 025 0.48a 0.45 0.47 0.168 2.445 0.017 026 0.23a 0.24 0.25 0.207 2.723 0.016 027 0.37a 0.38 0.37 0.189 2.441 0.019 028 0.24a 0.21 0.23 0.217 2.724 0.016 0

a from reference [6] and b from reference [7].

Table 2. Internal validation parameters of the MLR and PLS models (training set).

Model Ntraining RX2 r2

training q2L70

RMSEtr MAEtr CCCtr r2Scr q2

Scr

MLR_RR 21 - 0.917 0.748 0.029 0.025 0.957 0.198 -0.453MLR_RS 21 - 0.830 0.658 0.042 0.032 0.907 0.199 -0.422MLR_SR 21 - 0.863 0.743 0.037 0.029 0.926 0.149 -0.309MLR_SS 21 - 0.889 0.800 0.034 0.025 0.941 0.152 -0.340PLS_RR 21 0.726 0.976 (4)** 0.881 0.016 0.012 0.988 0.635 -0.510PLS_RS 21 0.701 0.965 (4)** 0.754 0.019 0.014 0.982 0.627 -0.557PLS_SR 21 0.702 0.972 (4)** 0.792 0.017 0.013 0.986 0.556 -0.571PLS_SS 20* 0.461 0.885 (2)** 0.656 0.031 0.026 0.939 0.324 -0.398

* Compound 5 was found as outlier and omitted from the final model; ** Number of components is given in parenthesis.

Table 3. External validation parameters of the MLR and PLS models (test set).Model Q2

F1 Q2F2 Q2

F3RMSEext MAEext CCCext

MLR_RR 0.811 0.808 0.833 0.041 0.037 0.900MLR_RS 0.814 0.811 0.836 0.041 0.033 0.901MLR_SR 0.713 0.708 0.747 0.051 0.044 0.856MLR_SS 0.787 0.783 0.812 0.044 0.039 0.896PLS_RR 0.777 0.773 0.803 0.045 0.039 0.912PLS_RS 0.727 0.723 0.759 0.050 0.040 0.902PLS_SR 0.716 0.711 0.749 0.051 0.042 0.898PLS_SS 0.554 0.529 0.514 0.050 0.056 0.799

Crisan, L., Iliescu, S., & Funar-Timofei, S.

Polímeros, 26(2), 129-136, 2016134

In the PLS modelling the terms having VIP values greater than 1 are the most relevant for explaining the dependent variable, and usually only these descriptors were interpreted. The descriptors showing the largest VIP values can simulate polymer flammability and are discussed below.

For all models higher values of the Randic shape index (-path/walk 4 and path/walk 5 - PW4 and PW5) are favourable for the flammability, while the MSD (Balaban mean square distance index) descriptor is unfavourable for flammability. They are topological descriptors obtained from molecular graph[15].

Another group of significant descriptors is the class of 2D autocorrelation descriptors, which are computed from molecular graph as the sum of products of atom weights of the terminal atoms of all the paths for the considered path length (the so called lag)[15]. The most important 2D autocorrelation descriptors involved in our model are the Geary parameters. The positive coefficients of GATS6m - Geary autocorrelation of lag 6 weighted by mass, increase the flame retardancy of RR, RS and SR series, while for SS dimers, the same effect was observed for descriptor GATS5v - Geary autocorrelation of lag 5 weighted by van der Waals volume.

The 3D-MoRSE descriptors provide 3D information from atomic coordinates using the same transform as in electron diffraction (which uses them to prepare theoretical scattering curves)[15]. For the RR and SR datasets, Mor13m- signal 13/weighted by mass, decrease the flame retardancy, for the RS dataset Mor15m- signal 15/weighted by mass is favourable for flammability, while for SS dataset these descriptors are insignificant.

Class of topological and frequency fingerprints descriptors are expressed as sum of topological distances between two elements or frequency of two atoms at a topological distance. Descriptors T(O..P) - the sum of topological distances between O..P, F07[O-S] – the frequency of O - S at topological distance 7, and F10[C-S] - the frequency of C - S at topological distance 10, with negative coefficients are unfavourable for flammability for RR, RS and SS datasets.

Three GETAWAY descriptors: one in the RR set: R5m- the R autocorrelation of lag 5/weighted by mass, and two in the SR set: HATS6v – the leverage-weighted

autocorrelation of lag 6/weighted by van der Waals volume and HATS6p- leverage-weighted autocorrelation of lag 6/weighted by polarizability increase the dimer flame retardancy.

Better fitting and predictivity results were obtained by PLS calculations compared to the MLR ones. From MLR and PLS models better statistical results were observed in case of the RR series. Therefore R chirality of phosphorous atom is significant for dimer flammability. The final selected structural descriptors included in the MLR_RR model have VIP values > 1 in the PLS_RR model: EEig09d, VIP = 1.358, CoeffCS = -0.0086 (±0.0022), MSD, VIP = 1.670, CoeffCS = -0.0096 (±0.0021), R2m, VIP = 1.058, CoeffCS = 0.0025 (±0.0022) and nP(=O)O2R, VIP = 0.994, CoeffCS = 0.0059 (±0.0030).

Compared to the MLR previously published monomer models[6], the statistical results for fitting are improved in case of MLR and PLS dimer models. Additional structural information which influences the flame retardancy was included in the final dimer MLR (e.g. the number of phosphonates) and PLS (e.g. 2D frequency fingerprints) models.

4. Conclusions

The MLR and PLS models developed for this series of dimer phosphoesters will be helpful to predict the LOI values of new untested compounds. Better statistical results and a more stable model to simulate polymer flammability were noticed in case of the RR dataset compared to the others, the presence of R chiral centre at the phosphorous atom being important for the dimer flammability. The mean square distance index and GETAWAY descriptors favour the dimer flammability, as well as increased number of phosphonates included in the dimer structure, as derived from both MLR and PLS methodologies. Better PLS fitting and predictivity results were obtained compared to the MLR ones for all datasets, except for the SS one.

Dimers including structures with R chiral centres gave more stable and predictive models compared to the previously published MLR monomer ones.

New structural information which influences the flame retardancy was included in the final MLR and PLS dimer models.

Table 4. Golbraikh and Tropsha criteria[29] calculated for external validation of the MLR and PLS models (test set).

Model 2testr

2 20

2r r

r− 2 '2

02

r rr− k k’ 2 '2

0 0r r−

MLR_RR 0.832 0.000 0.039 0.963 1.028 0.032MLR_RS 0.867 0.004 0.050 0.945 1.049 0.040MLR_SR 0.732 0.032 0.022 0.994 0.987 0.008MLR_SS 0.812 0.032 0.002 1.014 0.973 0.025PLS_RR 0.941 0.022 0.006 0.912 1.091 0.015PLS_RS 0.970 0.037 0.019 0.896 1.115 0.018PLS_SR 0.940 0.050 0.018 0.903 1.100 0.030PLS_SS 0.687 0.073 0.018 0.923 1.059 0.038

Structure-flammability relationship study of phosphoester dimers by MLR and PLS

Polímeros, 26(2), 129-136, 2016 135

5. Acknowledgements

This project was financially supported by Project 1.1 of the Institute of Chemistry Timisoara of the Romanian Academy. The authors are indebted to Chemaxon Ltd., OpenEye Ltd. and Prof. Paola Gramatica from The University of Insubria (Varese, Italy) for giving access to their software.

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Received: July 20, 2015 Revised: Sept. 23, 2015

Accepted: Nov. 06, 2015

http://dx.doi.org/10.1590/0104-1428.2323

SSSSSSSSSSSSSSSSSSSS

Polímeros, 26(2), 137-143, 2016 137

Influence of PLGA and PLGA-PEG on the dissolution profile of oxaliplatin

Emiliane Daher Pereira1, Renata Cerruti1, Edson Fernandes1, Luis Peña2, Vivian Saez1,2, José Carlos Pinto3, José Angel Ramón1,2,3, Geiza Esperandio Oliveira1,4 and Fernando Gomes de Souza Júnior1,5*

1Laboratório de Biopolímeros e Sensores, Instituto de Macromoléculas, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil

2Centro de Biomateriales, Universidad de La Habana, Havana, Cuba3Programa de Engenharia Química, Instituto Alberto Luiz Coimbra de Pós-graduação e

Pesquisa de Engenharia – COPPE, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil4Departamento de Química, Universidade Federal do Espírito Santo – UFES, Vitória, ES, Brazil

5Programa de Engenharia Civil, Instituto Alberto Luiz Coimbra de Pós-graduação e Pesquisa de Engenharia – COPPE, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil

*fgsj@ufrj.br

Sbstract

Oxaliplatin was inserted into polymeric matrices aiming to study the interaction of this drug with these polymers and its capability to diffuse to the environment. Tested polymers were: (1) polyethylene glycol (PEG), (2) poly(lactic-co-glycolic acid) (PLGA), and (3) a copolymer of them (PLGA-PEG). The latter two were synthesized by us using polycondensation in bulk. Oxaliplatin was included in the matrices by the melt mixing process followed by casting. Fourier tran sform infrared spectroscopy (FTIR), proton nuclear magnetic resonance (1H-NMR) and X-ray diffraction (DRX) studies of the polymers were performed proving the obtaining of the desired materials. In addition, the interaction between drug and matrices and the release profile of the oxaliplatin from these matrices were analyzed. Among them, PEG did not control the oxaliplatin release. In turn, PLGA and PLGA-PEG present drug release profiles quite similar. Oxaliplatin was completely released from PLGA and PLGA-PEG in 5 hours, by a relaxation mechanism. There was no evidence of oxaliplatin interaction with the different polymers. In addition, as the PEG improves the biocompatibility and biomasking, obtained results prove the obtaining of a drug release system, which allowed the total use of the drug improving the cancer treatment and even the welfare of the patients.

Keywords: oxaliplatin, drug delivery, biodegradable polymer, PLGA-PEG, block copolymer.

1. Introduction

More than 10 million people are diagnosed with cancer annually, making this disease a leading cause of death. Furthermore, new projections allow predicting that by 2020, there will be 15 million new cases of cancer every year[1,2]. Nowadays, the third most common cancer worldwide is the colorectal cancer and this is also the fourth most common cause of death[3-6].

Oxaliplatin is a third-generation platinum-based chemotherapeutic agent, which presents a remarkable success on the cancer treatment, as the other platinum-based drugs, such as cisplatin and carboplatin[7]. Specifically, oxaliplatin is a organoplatinum compound with significant activity against advanced or metastatic digestive tumors, mainly colorectal ones[8]. Therefore, oxaliplatin is considered one of the first-line drugs in the treatment of advanced colorectal cancer[9-12]. In addition, oxaliplatin does not induce nephrotoxicity to the same extent as these other platinum-based drugs[13]. On the other hand, recent reports criticize the oxaliplatin use due to some cases of peripheral neuropathy caused by cumulative doses of this drug[13,14].

Drugs entrapped into biodegradable and biocompatible polymeric matrices have been successfully employed for the

controlled release of drugs[15-19]. Potentially, these systems are able to maintain drug concentrations within therapeutic ranges even along days. Consequently, it may diminish side effects caused by high concentrations and repeated administrations, and improve patient compliance as compared to conventional regimens[20]. Poly(lactic-co-glycolic acid) (PLGA) copolymers[21] have been extensively used in drug delivery systems (DDS) as alternative to improve conventional formulations[21-23]. These DDSs allow extended and controlled releases of drugs dispersed into the polymer[24]. However, PLGA hydrophobicity may be a problem for some drugs. Incomplete release is usual if drugs have interactions with them[25]. The attaching of PEG chains to PLGA may solve these situations[26,27]. The inclusion of PEG into material could also be convenient for improving biocompatibility of the systems[28]. However, high solubility of PEG could accelerate the drug release rate which in some cases it is an undesirable effect.

In this paper, the effect of the hydrophilicity of the polymeric matrix on the release profiles of oxaliplatin was studied. In addition, possible interactions between the drug and the matrices were investigated. As expected, PEG did not

Pereira, E. D., Cerruti, R., Fernandes, E., Peña, L., Saez, V., Pinto, J. C., Ramón, J. A., Oliveira, G. E., & Souza, F. G., Jr.

Polímeros, 26(2), 137-143, 2016138

act as a release controlling matrix for oxaliplatin. However it is possible to exploit the biocompatibility provided by the PEG as part of PLGA-PEG copolymer because its presence does not change the drug release profile. Oxaliplatin was completely released from both PLGA and PLGA-PEG in five hours.

2. Experimental Part

2.1 Materials

PEG (Mn 6000) was purchased from SIGMA-ALDRICH, Brazil. The oxaliplatin was gently provided by Center for Research and Development of Medicine, Havana, Cuba. All others chemical used were purchased from VETEC, Brazil. All chemical materials were used as received.

2.2 Synthesis of block copolymers

The PLGA block copolymer was synthesized by polycondensation in bulk from lactic acid (50 ml) and glycolic acid (62 ml) in equimolar ratio (1:1). Sulfuric acid (98% v/v, 0.25mL) was added as catalyst, in a closed system under nitrogen atmosphere and slight vacuum. Temperature was kept at 140°C under magnetic stirring for 10 h[29]. The PLGA-PEG block copolymer was synthesized following the same procedure, with addition of commercial PEG (5g)[27].

2.3 Preparation of polymer/drug systems

The three polymers (PEG; PLGA; PLGA-PEG) were molten separately at 120°C. This temperature was chosen since it is considerably lower than the degradation temperature of the drug which is 400°C according to the results of the TGA analysis. Oxaliplatin was incorporated in the three polymers in a 1% w/w concentration by melting at 120°C inside a 50mL reactor using mechanical stirring and dry N2(g) flow. After the drug incorporation, melted mixture was casted, producing disks (diameter: 1.0 cm; height: 0.5 cm). The DDS prepared were named as PEGO, PLGAO and PLGA-PEGO, respectively.

2.4 Materials characterization

Hydrogen nuclear magnetic resonance (1H-NMR) and Fourier transform infrared (FT-IR) were used to study the composition of the copolymers. Samples were dissolved in CDCl3 and their 1H NMR spectra were recorded on a Varian equipment model Oxford 300. FTIR spectra were obtained from neat films cast from the chloroform sample solutions on KBr tablets, with a Varian equipment model 3100 FTIR Excalibur Series, using a resolution of 4 cm-1 and 20 scans from 4000 to 400cm-1.

X-Ray Diffraction (XRD) was also used to characterize the copolymers, the samples prepared and the drug. The equipment used was a Rigaku Miniflex X-ray diffractometer in a 2θ range from 2° to 80° by the method FT (fixed time). The steps were equal to 0.05°/s, using a tube voltage and current equal to 30 kV and 15 mA, respectively. The radiation used was CuKα = 1.5418 Å.

2.5 Presence of oxaliplatin in matrices

The technique of least squares was used to compare the FTIR data obtained from the DDS containing oxaliplatin with the one of matrices without the drug. Specifically, the coefficient of determination (R2) and the root mean square error (RMSE) were used as indicatives of the presence of oxaliplatin into the matrices. This test was reported by our group elsewhere[30].

2.6 Dissolution test of oxaliplatin

For each test, one gram of PEGO, PLGAO and PLGA-PEGO was studied using the USP Apparatus I at 75 rpm. First, the released amount of the drug was analyzed in HCl 0.1 M (900 ml) by one hour. Soon afterwards, this solution was substituted by the same volume of phosphate buffer pH 7.8 and drug elution was determined by other 4 hours. Samples were collected at 0.25h; 0.5h; 1h; 1.25h; 1.5h; 2h; 3h; 4h and 5h. All these tests were performed in triplicate.

Ultraviolet-visible spectrophotometry (UV-Vis) was used to quantify the oxaliplatin in the samples. The measurements were performed at 260 nm using a monochromator Biospectro spectrophotometer model SP-220. The spectra of the solutions with known concentrations of oxaliplatin were collected and used to set up an analytical curve.

3. Results and Discussion

Figure 1 shows the 1H-NMR spectrum of the PLGA copolymer. The peak at 1.5 ppm is associated with the repetitive methyl groups of the lactic acid. The peaks at 5.2 ppm and 4.8 ppm are related to the CH-CH3 from lactic acid and CH-H from glycolic acid, respectively, confirming the synthesis of the copolymer[27,29]. The molar amount of each comer of the copolymer was calculated using the relationship among the peak areas and the number of H atoms. The synthesized copolymer presented 46.6mol% of their molar mass related to the glycolate comer, while the lactate comer corresponded to 53.4mol% of the copolymer.

Figure 2 shows the 1H-NMR spectrum of the PLGA-PEG copolymer. In this spectrum, same peaks of PLGA are present

Figure 1. 1H-NMR spectrum of the PLGA copolymer.

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Polímeros, 26(2), 137-143, 2016 139

and the new peak at 3.64 ppm is assigned to the methylene groups present in PEG[27,31-33]. Again, the peak areas were used to estimate the molar amount of PEG in the system PLGA-PEG. Calculated value is around 5.9mol%, while the molar mass related to the glycolate and lactate comers corresponds to 45.7mol% and 48.4mol%, respectively.

FTIR spectra of PEG, PLGA and PLGA-PEG are shown in Figure 3. The FTIR spectrum shown in Figure 3a corresponds to PEG. The smooth wide band around 3500 cm-1 is related to the oxygen atom able to form hydrogen bonds. The characteristic band at 2888 cm-1 is attributed to the stretching of CH, CH2 and CH3 groups. The characteristic bands at 1464 and 1343 cm-1 are associated with the C-H bending. The stretching of C-C-O group appears at 1280 and 1100 cm-1. The later characteristic band is conjugated with the C-O-C group. The harmonic bands of C-C-O group appear around 950 and 840 cm-1. The spectrum shown in Figure 3b corresponds to PLGA copolymer, again is possible to see a smooth wide band around 3500 cm-1 related to the oxygen capability to form hydrogen bonds. The small doublet around 3000 cm-1 is related to CH, CH2 and CH3 groups. The characteristic peak at 1752 cm-1 is attributed to the stretching of the C=O. While the characteristic band at 1182 cm-1 can be assigned to the ether group. The characteristic bands at 1130 cm-1 and 1452 cm-1 are attributed to C-O-C group and C-H bond of the methyl group, respectively[32,34-37]. The Figure 3c shows the PLGA-PEG spectrum. The spectrum profile is very similar to the PLGA one, which can be explained by the small amount of PEG in the system, around 5.9mol%. As showed by the spectra, the PLGA characteristic bands are stronger than PEG, since the former corresponds to more than 90% of the block copolymer. Therefore, the PEG presence can be barely noticed by the small signal of the C-C-O harmonics, placed around 950 and 840 cm-1.

The FTIR spectra of oxaliplatin and the samples PEGO, PLGAO, PLGA-PEGO are showed in Figure 4. In the spectrum of oxaliplatin (Figure 4a), the peaks at 3264 and 3509 cm–1 confirm the presence of an NH stretch, and peak at 812 cm–1 shows N-H bending; C=O stretch was observed at 1707 cm–1. In spectra of matrices combined with oxaliplatin (Figure 4b, 4c and 4d), signals corresponding to drug were not observed, maybe due to its low concentration in the mixtures[38].

Due to small differences among the FTIR spectra of the matrices in comparison to the ones containing oxaliplatin, the RMSE statistical tool was used to prove the presence of the drug in the materials. These differences can be expressed using appropriate statistical tools. Among them, the root mean squared error (RMSE) is very useful to study the misfit between experimental data and model[30,39]. Therefore, RMSE was calculated through linear regressions between absorbances of the matrix and analogous filled with oxaliplatin, using the least squares approach[40]. The obtained results are shown in Figure 5. As reference, absorbance of the PLA was plotted as a function of its own absorbance and RMSE is, obviously, equal to zero. In turn, absorbances of PEGOxPEG, PLGAOxPLGA and PLGA-PEGOxPLGA-PEG produced RMSE values equal to 4.79, 3.27 and 3.11, respectively. In spite observed differences, probably produced by the chemical differences among tested materials, all of these

results were obtained with p<0.05, indicating that matrices were successfully loaded with oxaliplatin.

Figure 6 shows the XRD analysis for polymers and DDS. The diffractogram of PEG (Figure 6a) exhibits two peaks indicative of crystallinity: 19.2°(2θ) and 23.4°[41]. The diffractogram of pure oxaliplatin (see Figure 6d) exhibits many characteristic peaks due to its crystalline nature[38]. The diffractograms of PLGAO and PLGA-PEGO (see Figure 6f and g), respectively) show amorphous halos of the copolymer and the characteristic peaks of the drug. In turn, the XRD of the PLGA-PEG (Figure 6c) also shows only amorphous halos due the low proportion of the PEG in the copolymer (3.5%) and PLGA (Figure 6 (b)) presents only amorphous halos which are in agreement with the literature[34].

Figure 4. FTIR spectra of oxaliplatin (a), PEGO (b), PLGAO (c) and PLGA-PEGO (d).

Figure 2. 1H-NMR spectrum of the PLGA-PEG copolymer.

Figure 3. FTIR spectra of PEG (a), PLGA (b) and PLGA-PEG (c).

Pereira, E. D., Cerruti, R., Fernandes, E., Peña, L., Saez, V., Pinto, J. C., Ramón, J. A., Oliveira, G. E., & Souza, F. G., Jr.

Polímeros, 26(2), 137-143, 2016140

Dissolution tests of oxaliplatin-loaded DDS were monitored by UV-Vis analysis. Results are shown in Figure 7. Total drug release from PEG matrix (Figure 7a)

took place in one hour due to its high hydrophilic character. The PLGA (Figure 7b) and PLGA-PEG (Figure 7c) matrices control effectively the release rate of oxaliplatin, which was completely delivered in five hours. At the end of study disks were completely degraded. The experimental data corresponding to these samples was fitted using the model presented in the Equation 1:

( )1 2

/12x x0

A ARAO = + Adx+e −

− (1)

In Equation 1, RAO is the released amount of oxaliplatin, A1 is the lower RAO limit, A2 is the higher RAO limit, x0 is the inflexion (half amplitude) point and dx is the width.

Results are showed in Table 1. All the parameters are statistically equal, confirming that oxaliplatin is released from PLGA and PLGA-PEG matrices in the same way.

The knowledge about the drug release mechanism and release rates is very important for programming the characteristics of the systems. Consequently, mathematical modeling of drug release processes from polymeric matrices plays an important role to predict the behavior of the system and to determine the structural parameters of the polymer that affect and control the drug release profiles. The drug release can occur through a pure Fickian diffusion, by

Figure 5. Comparison between normalized absorbances of oxaliplatin-loaded systems and non-loaded copolymers. PLGA vs PLGA (a), PEGO vs PEG (a), PLGAO vs PLGA (b) and PLGA-PEGO vs PLGA-PEG (d).

Figure 6. X-ray diffraction patterns of PEG (a), PLGA (b), PLGA-PEG (c), PEGO (d), PLGAO (e), PLGA-PEGO (f) and pure oxaliplatin (g).

Influence of PLGA and PLGA-PEG on the dissolution profile of oxaliplatin

Polímeros, 26(2), 137-143, 2016 141

processes of relaxation and sometimes a combination of these two mechanisms takes place[42]. When the semi-empirical equation of Ritger-Peppas, [43], is used for fitting the experimental data, the diffusional exponent (n) indicates the mechanism for drug transport. An exponent value for cylindrical geometric n = 0.45 indicates Fickian diffusion, n within 0.45-0.89 indicates the anomalous transport, n = 0.89 indicates case II transport and n more than 0.89 indicates super case II transport. An anomalous transport (0.45<n<0.89) is referred to a drug transport by diffusion and relaxation of the polymer chains[43].

The polymeric materials that controlled the release of oxaliplatin (PLGA and PLGA-PEG) have a value for Tg lower than the temperature for running the dissolution tests[44]. In that sense it is reasonable to suppose that drug transport could occur by a combination of both: diffusion and relaxation. Thus, for obtaining a preliminary idea of the release mechanism for oxaliplatin included in these disks, experimental data was fitted to theoretical release profiles given by a Peppas-Sahlin equation (Equation 2), which combines the Fickian diffusion and non-Fickian movement of drug molecules through polymer chains[45].

2m mt1 2

M = k t +k tM∞

(2)

In Equation 2, k1 and k2 are Fickian and relaxational contribution respectively[46]. They allow calculating the approximate contributions of the diffusional and relaxational mechanisms.

Experimental modeling was performed using the data from drug released until one hour, which corresponded to 60% of the total drug contained in DDS. The exponent m was assumed equal to 0.44 due to geometry of the samples, which were disks. For PLGAO, k1 was -0.77 ± 0.01 and k2 was 0.91 ± 0.01; for PLGA-PEGO, k1 was -0.93 ± 0.02 and k2 was 1.05 ± 0.02 (R2 in both cases was greater than 0.99). A negative value of k1 was obtained in both cases, which can be considered as an insignificant effect of Fickian diffusion on drug release compared to the relaxation process. This result can be related to Tg of the polymers which are

lower than 37ºC. In that sense when disks are placed into release fluid the combination of heating and the swelling in water could favor the polymer change from glassy to rubbery state. Consequently, in this case, relaxation of the polymer is the driving force for the release of oxaliplatin.

4. Conclusions

This work presented the release of oxaliplatin from PEG, PLGA and PLGA-PEG matrices. PLGA and PLGA-PEG were obtained by polycondensation and the synthesis of products was confirmed by 1H-NMR, FTIR and XRD. The release studies, monitored by ultraviolet analysis, showed that the PEG matrix cannot sustain the release for more than one hour, maybe due to the high solubility of this polymer. In contrast, the PLGA and PLGA-PEG copolymers release oxaliplatin during 5 hours in a similar way. So, the modification of PLGA using PEG did not modify oxaliplatin release. In addition, the characterization of materials by FTIR and the complete release of oxaliplatin from DDS allow believing that there are not interactions between drug and copolymers while the release mechanism seems to be due to relaxation of the polymer chains. Therefore, the PLGA-PEG copolymers could be preferred for designing controlled release systems for oxaliplatin with enhanced biomasking and biocompatibility.

5. Acknowledgements

The au tho r s t hank t o Conse lho Nac iona l de Desenvolvimento Cient í f ico e Tecnológico (CNPq-474940/2012-8 and 550030/2013-1), Coordenação de Aperfeiçoamento de Pessoal de Nível Superior (CAPES and CAPES-NANOBIOTEC), Financiadora de Estudos e Projetos (FINEP PRESAL Ref.1889/10) and Fundação Carlos Chagas Filho de Amparo à Pesquisa do Estado do Rio de Janeiro (FAPERJ) for the financial support and scholarships.

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Table 1. Parameters of the dissolution curves of oxaliplatin released from PLGA and PLGA-PEG matrices.Sample A1 A2 dx x0 R2

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Figure 7. Oxaliplatin released along experimental time (first hour HCl 0.1M and last 4 hours with phosphate buffer): system PEGO (a); system PLGAO (b) and system PLGA-PEGO (c).

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Received: Aug. 14, 2015 Revised: Oct. 07, 2015

Accepted: Nov. 09, 2015

http://dx.doi.org/10.1590/0104-1428.2294

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Polímeros, 26(2), 144-151, 2016144

Biopolymer production using fungus Mucor racemosus Fresenius and glycerol as substrate

Thaíssa Rodrigues Araújo1, Carmen Lúcia de Oliveira Petkowicz2, Vicelma Luiz Cardoso1, Ubirajara Coutinho Filho1 and Patrícia Angélica Vieira1*

1Faculdade de Engenharia Química, Universidade Federal de Uberlândia – UFU, Uberlândia, MG, Brazil2Departamento de Bioquímica, ACF Centro Politécnico, Universidade Federal do Paraná – UFPR, Curitiba, PR, Brazil

*patriciavieira@feq.ufu.br

Sbstract

This study evaluated extracellular production of biopolymer using fungus Mucor racemosus Fresenius and glycerol as a carbon source. Initially employing conical flasks of 500 mL containing 100 mL of cultive medium with 0.18 ± 0.03 g.L–1 of microorganisms, the results showed that the best conditions of the variables studied were: initial concentration of glycerol 50 g.L–1, fermentation time of 96 h, inoculum cultivation time of 120 h, and aeration in two stages–the first 24 hours without aeration and 72 hours fermentation with aeration of 2 vvm and 2 g.L–1 of yeast extract. The experiments conducted in a Biostat B fermenter with a 2.0 L capacity that contained 1.0 L of medium showed production of 16.35 g.L–1 gum formed and 75% glycerol consumption. These conditions produced a biopolymer with the molecular weight and total sugar content of 4.607×106 g.mol–1 (Da) and 89.5%, respectively.

Keywords: biopolymer, Mucor racemosus Fresenius, glycerol, aeration, fermentation, productivity.

1. Introduction

The excess of glycerol in the market, generated by the production of biodiesel, is a problem that has been discussed over the years. Researchers are seeking an alternative for the bioconversion of this by-product. According to Moralejo-Garate et al.[1] the crude glycerol is a by-product of the biodiesel industry and a potentially good substrate for the production of biopolymers.

Biopolymers are complex chains of polysaccharides of microbial origin, synthesized by bacteria, fungi and yeasts, and are also known as gums due to their ability to form viscous solutions and gels in aqueous media. The biopolymers produced by fungi have not been adequately explored, and only some of them have been produced on an industrial scale. However, several of these biopolymers have attracted attention due to their physico-chemical and rheological properties, and have found a wide range of applications, including use in pharmaceutical therapy for action-tumors, anti-virals and anti-inflammatories[2].

Biopolymers have unique chemical and physical properties that are superior to those of traditional polysaccharides, such as higher viscosity and gelling power, compatibility with a wide variety of salts in a large range of pH and temperature, stable in high ion concentrations, high solubility water, and also synergistic action with other polysaccharides[3,4].

The increasing interest in biopolymers compared to traditional polymers is due to its scarcity and high oil prices in addition to the environmental impact that is determined by its extraction, refining and difficult biological degradation[5].

The species of the genus Mucor has been reported in the literature as a potential producer of high value bioproducts. According to Alves et al.[6] the mucor species genera are

a large group of fungi with potential biotechnological importance, which is responsible for the production of industrial enzymes. The high production capacity of lipase, protease and phytase by Mucor racemosus[6-7] are mentioned in the literature. Furthermore, the strains Mucor racemosus are cited in the literature as producing various bioproducts, such as the production of chitosan, which is extracted from the mycelium of this strain in various growth stages[8]. In the literature, there are few papers based on the use of this fungus for the production of biopolymers. Nevertheless, none of them report the production of the biopolymer applying glycerol as substrate.

The goal of this study was to evaluate the biopolymer production employing the fungus Mucor racemosus Fresenius and glycerol as substrate and to select the best operational conditions among those studied. The subsequent objective was to evaluate the ability of the fungus to produce biopolymer using a 2 L capacity Biostat B fermenter.

2. Materials and Methods

2.1 Microorganisms, medium culture for maintenance and production process2.1.1 Microorganisms

In this work we used a fungus isolated from a central area in Brazil (State of Minas Gerais) identified by the Foundation André Tosello for Research and Technology (Campinas, SP, Brazil) and Laboratory Exame (Uberlândia, MG, Brazil) as Mucor racemosus Fresenius. This microorganism was selected among other fungi in preliminary trials as a potential producer of biopolymer.

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2.1.2 Maintenance of microorganisms: cultivation of inoculums

The fungus was grown in Petri dishes containing Czapeck medium (pH=6) with the following composition in g.L–1: 20.0 glucose, 2.0 sodium nitrate, 1.0 potassium phosphate dibasic, 0.5 magnesium sulfate, 0.5 potassium chloride, 0.01 ferrous sulfate, 20.0 agar-agar. The medium was sterilized at 110 °C and 1 atm for 20 minutes. The incubation was performed at 28 ± 1 °C for 120 h (initial standard time).

2.1.3 Biopolymer production

The culture medium used in all experiments, obtained in preliminary tests, showed the following composition in g.L–1: 0.5 K2HPO4, 1.4 KH2PO4, 1.0 NH4NO3, 1.0 MgSO4.7H2O, 0.02 CaCl2 .H2O, 0.03 MnSO4.H2O, 2.0 yeast extract and varying concentrations of glycerol in accordance with each assay. The medium had its pH adjusted to 6.5 with subsequent sterilization at 110 °C and 1 atm for 20 minutes.

2.2 Carbon source

Commercial glycerol (Vetec, Brazil) with a purity of 99.5% was used as substrate.

2.3 Preliminary tests: evaluation of process variable

All experiments were performed in conical flasks of 500 mL containing 100 mL of culture medium, and glycerol concentrations varied according to each experiment. The amount of inoculum used corresponded to 0.18 ± 0.03 g.L–1.

The variables evaluated in chronological sequence were: initial concentration of glycerol (without aeration process), fermentation time, cultivation time of the inoculum (inoculum age), aeration flow, aeration mode and the effect of yeast extract in the biopolymer production.

All experiments were conducted twice and triplicate.Determination of glycerol concentration, biopolymer

production time and inoculum cultivation timeThe experiments for evaluating the optimal concentration

of glycerol, biopolymer production time and inoculum cultivation time were performed without additional aeration.

The response variables monitored in the experiments were the following: concentration of the biopolymer, glycerol consumption, productivity (PP) and the product yield constant (g-product/g-substrate) (YP/S). All assays were performed under agitation in an oscillatory shaker (New Brunswick) at 150 rpm and 30.0 ± 2 °C. The productivity and product yield constant is given by Equations 1 and 2, respectively:

0mP

fp

P PPt−

= (1)

0

0

mPS

P PYS S−

=−

(2)

Where, PP: productivity;P0: initial product concentration;Pm: maximum product concentration (final);tfp: product formation time;S0: initial substrate concentration (glycerol);

S: final substrate concentration (glycerol);YP/S: product yield constant (g-product/g-substrate).

The glycerol initial concentrations tested were 25, 50, 100 and 125 g.L–1. For the fermentation time, 48, 72, 96, 120, 144 and 168 hours were tested. In order to verify the effect of inoculum cultivation time (age of inoculum) in the biopolymer production process, the following times were evaluated: 48, 72, 96 and 120 hours.

2.3.1 Evaluation of aeration mode in the production process of biopolymer

Initially, a preliminary test was performed with continuous aeration of 3 vvm. In this operation (3 vvm) was observed the decrease in the medium volume in the reactor (the liquid medium was entrained in air). And the production process of gum was impaired. Subsequently, three additional tests were performed. The first test was performed without aeration; the second test with continuous aeration of 2 vvm. The third test was carried out without aeration in the first 24 hours and with continuous aeration of 2 vvm in 72 hours of assay.

2.3.2 Evaluation of the effect of yeast extract addition on biopolymer production

After selecting the initial concentration of glycerol (without aeration process), fermentation time, cultivation time of the inoculum (inoculum age), aeration flow, and aeration mode, the effect of yeast extract concentration in the production of the biopolymer was valued.

To evaluate the effect of yeast extract (YE) in biopolymer production, different concentrations were tested: 0; 0.5; 1; 2; 2.5, 3.0 and 4.0 g.L–1. It should be emphasized that a trial was also carried out using only yeast extract as the source of carbon at concentrations of 1, 2, 3 and 4 g.L–1, i.e., without glycerol.

2.4 Biopolymer production in Biostat B fermenter

After optimization of variables (inoculum cultivation time, glycerol concentration, level of aeration, fermentation time and yeast extract concentration) kinetic assays were performed in Biostat B fermenter to evaluate the performance of the fungus in the extracellular production of biopolymer. The experiment was performed using a Biostat B fermenter with a 2.0 L capacity that contained 1.0 L of medium, with the first 24 h without aeration and with 72 h with aeration of 2 vvm. The above procedure was used, because preliminary tests showed that after 24 h fermentation had visible change in the viscosity of the fermentation broth. This fact show that gum production occurred after that time. Therefore, it was tested the addition of aeration after 24 hours of fermentation.

The process was evaluated by monitoring the amount of formed gum, glycerol consumption and cell concentration. For the characterization of the biopolymer produced, the following analyses were performed: rheologic behavior evaluation, monosaccharide composition and molecular parameters by steric exclusion chromatography.

2.5 Analytical determinations2.5.1 Recovery and purification of biopolymer

The fermented broth was diluted 1:1 with deionized water and centrifuged in a Beckman Coulter Avanti J-25 centrifuge at 18,900×g for 40 min to remove cells.

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This procedure of dilution was necessary, because the medium was very viscous. This facilitated the withdrawal of the fermented broth cells. This procedure was the same adopted by Faria et al[9]. The supernatant was filtered and treated with a saturated solution of KCl, and the polymer was recovered by precipitation with ethanol. Finally, the product was dried under vacuum system at 30 ± 1 °C[10].

2.5.2 Analysis of glycerol concentration

Glycerol concentration in the cell-free supernatant was determined by liquid chromatography (HPLC) with an Aminex HPX-87H column (Bio Rad), coupled to a refractometer. The analysis was performed at 50 °C, with sulphuric acid (H2SO4 0.01 N) as eluent, at a flow rate of 0.6 mL/min[11].

2.5.3 Determination of biomass

After fermentation, the fermented culture medium was diluted 1:1 in water, and the cells were separated by centrifugation Beckman Coulter Avanti J-25-18900 g for 30 minutes. The precipitate (cells) was washed three times in distilled water for complete removal of medium components and metabolites[10].

2.5.4 Evaluation of rheological behavior of the biopolymer

The rheology of gum was obtained by preparing a polymer solution of 0.5%, 0.75% and 1% under magnetic stirring for approximately 10 h. The gum used to prepare this solution was dried at room temperature, and then crushed and hydrated[10,12].The rheology was measured using a Brookfield rheometer RVDVIII coupled with a water bath mark Brookfield, model TC-502P, using the small sample adapter, spindle 18. Following the template Ostwald Waele or power-law, data on shear stress, measured from shear rates, were used. The units of measurement were apparent viscosity in Pa.s and s–1 to shear rate.

2.5.5 Determination of monosaccharide composition and molecular parameters

To determine the monosaccharide composition of the fractions in terms of neutral sugar samples were hydrolyzed with 2 M trifluoroacetic acid in two conditions: for 8 h at 100 °C or 120 °C for 2 h. For each hydrolysis condition, four replications were performed. Upon completion of the hydrolysis, the excess of acid was removed by evaporation[13]. After total acid hydrolysis, the monosaccharides were solubilized in about 5mL of distilled water and reduced by adding approximately 10 mg of sodium borohydride for 16 h at 4 °C. Later, strongly acidic cation exchange resin was added to remove the Na+ ions.

The solutions were filtered and the solvent evaporated in vacuo. Methanol (1 mL) was added to remove boric acid and borate of methyl formed was evaporated in vacuo. This process was repeated three times. The formed alditols were acetylated by adding 0.5 mL of acetic anhydride and 0.5 mL of pyridine in sealed tubes, standing for 12 h at room temperature. The reaction was stopped by the addition of ice, and then the extraction of ethyl alditols was performed by the addition of chloroform and subsequent elimination of pyridine in successive treatments with 5% copper sulphate and distilled water. After evaporation of the solvent, the acetates and the alditols were subjected to gas-liquid chromatography-mass spectrometry (GC-MS) in order to

determine the composition of the neutral monosaccharides. The monosaccharide composition was calculated as the average of eight replications.

2.5.6 Determination of the content of acid monosaccharides (uronic acids)

Dosing of uronic acid concentration was accomplished by the method of Blumenkrantz and Asboe-Hansen[14], with the standard solution galacturonic acid in concentrations of 10-100 μg. mL–1 and reading of 520 nm. A total of eight replications were carried out.

2.5.7 Identification of uronic acids

For the identification of the uronic acids, hydrolysed and free acid samples were filtered through cellulose acetate membranes with pores of 0.22 μm, and for anion exchange chromatography using a Dionex ICS-5000 and CarboPac PA20 column, according to the methodology described by Nagel et al.[15].

2.5.8 Liquid-gas chromatography mass spectrometry (LGC-MS)

The analyses were performed on a Varian 3300 gas chromatograph coupled with a mass spectrometer FINNIGAN-MAT, with injector the 50 °C and “ramp” of 40 °C per minute to 220 °C, equipped with a fused silica capillary column (30 m × 0.25 mm d.i) coated with DB-225, and helium as the carrier gas (1 mL.min–1).

2.5.9 Sample analysis by steric exclusion chromatography coupled with the detector of laser light of multi-angle and differential detector of index of the refraction (HPSEC-MALLS/RI)

Analyses were performed on a device consisting of an HPLC pump (Waters 515), injector, four columns of Ultrahydrogel–120, 250, 500 and 2,000–with limits of exclusion 5.103, 8.104, 4.105 and 7.106 respectively, DAWN DSP Light Scattering (Wyatt Technology), and an index detector of differential refractive model 2410 (Waters). The eluent used was a 0.1 M NaNO2 solution containing 200 ppm NaNO3.

The samples were solubilized at a concentration of 1 mg.mL–1 in the solution of the eluent. Before analysis, the samples were filtered through cellulose acetate membranes with pores of 0.22 micrometers. The average molar mass was determined by light scattering method.

2.5.10 Thin layer chromatography (TLC)

The monosaccharides were identified by Thin-Layer chromatography using a silica gel plate (20×20 cm) as stationary phase, previously activated in a greenhouse at 100 °C for 1 h. Monosaccharides were visualized using orcinol-H2SO4 (specific for carbohydrates) and then with ninhydrin (for amino group)[16].

3. Results and Discussion

3.1 Experiments employing stirred conical flasks

Figure 1 shows the influence of the initial concentration of glycerol in the formation of the biopolymer.

It can be observed that the glycerol concentration that promoted the higher formation of the biopolymer was 50 g.L–1. In this concentration, the content of gum

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was 2 g.L–1 with productivity of 0.012 g.L–1 h–1, product yield constant (YP/S) of 0.104 g of biopolymer/g glycerol and glycerol consumption of 38.4%. Furthermore, it can be seen by these results that higher concentrations of glycerol at 100 and 125 g.L–1 inhibited the production of gum. Nevertheless, at such concentrations, consumption of substrate was superior, achieving 42.75% and 46.12%, respectively. This suggests that consumed glycerol was used mainly for further purposes, besides the production of the biopolymer. The final concentration of cells for assays with glycerol concentration of 25, 50, 100 and 125 g.L–1 were: 0.34, 0.67, 1.01 e 1.21 g.L–1, respectively. This fact confirms the earlier hypothesis.

Antônio et al.[17], in their study of biopolymer production using bacteria Gluconoacetobacter hansenii and 25 g.L–1 of glycerol as substrate, obtained 6.9 g.L–1 of biopolymer after 14 days of fermentation. Comparing this result with what was found by the present work, it can be seen that despite a lower result presented in this study (2 g.L–1 gum and 50 g.L–1 glycerol in (7 days) 168 h of process), the studied microorganism showed a potential gum production capacity, since at this stage no further optimization of production conditions was carried out. Therefore, the concentration of 50 g.L–1 glycerol was chosen to be used in the following stages of this work.

Figure 2 shows the results of the amount of the produced gum and glycerol consumption in assays performed on fermentation times of 48, 72, 96, 120, 144 and 168 h.

Figure 2 shows that the maximum production of gum was obtained in the fermentation time of 96 h producing 2.96 g.L–1 of gum, productivity of 0.031 g.L–1 h–1, and product yield constant of YP/S of approximately 0.10 g of biopolymer/g of glycerol.

At 120, 144 and 168 h a decrease in the amount of gum formed was observed, corresponding to the values of 2.87, 2.26 and 2 g.L-1, respectively. This behavior suggests that enzymes produced by the own microorganism may have degraded the gum. This assumption may be confirmed by the consumption of glycerol according to the fermentation times illustrated in Figure 2, showing that glycerol consumption was enhanced despite the small amount of formed gum. This suggests that the gum begins to be degraded and the glycerol is consumed for the formation of other by-products. Based on these results, the fermentation time chosen to attain an increased amount of biopolymer production was 96 h.

The influence of fermentation time in producing gum is reported by some authors who argue that a longer fermentation time promotes higher productivity[18].

Figure 3 shows the results of the amount of formed gum and glycerol consumption using inoculum of different ages (incubation times). This study has become necessary since the incubation time is associated with the age of the inoculum, which can directly influence the process of fermentation, product generation and gum formation.

Figure 3 shows that the higher time of incubation, the higher the content of biopolymer produced up to 120 h. The use of an inoculum incubated for 120 h allowed a production of 3 g.L–1 (of fermentation time of 96 h), productivity of 0.031g.L–1 h–1, product yield constant of approximately, 0.100 g of biopolymer/g of glycerol and consumption of

glycerol of 33.9%. By applying an inoculum of 144 h, there was an increase of only 0.2 g.L–1 of gum formed, which is not statistically different. Thus, the age of the inoculum culture adopted for the development of this work was 120 h.

Figure 1. Amount biopolymer formed (■) and glycerol consumption (○) for different concentrations initial of glycerol substrate (25, 50, 100 and 125 g.L–1).

Figure 2. Study of the effect of fermentation time on the amount of formed gum (■) and glycerol consumption (○).

Figure 3. Effect of the incubation time of the inoculum about on the amount of formed gum (■) and glycerol consumption (○).

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In preliminary tests, in order to verify the effect of the aeration level values on the system of production of biopolymer, aeration of 3.0 vvm was tested along with other pre-established conditions. In this aeration maximum production of 4.04 g.L–1 gum and productivity of 0.084 g.L–1 were attained in 48 h. These results demonstrated that the aeration is an important variable in the gum formation by this fungus. Furthermore, it was observed that the air flow rate used was high for the 96 hours of processing, since drying of the culture medium occurred (the liquid medium was entrained in air). Another important point, checked with the aeration, was an increased productivity, which increased from 0.031 g.L-1 h-1 in the 96 h process to 0.084 g.L–1h–1 in 48 h. Therefore, aeration proved to be an important parameter to be explored and evaluated in different ways.

Therefore, in order to determine the best condition of aeration on the production of biopolymer, three experiments were performed, each one using 50 g.L–1 glycerol, cultivation time (inoculum age) of 120 hours, room temperature 30 ± 2 °C and fermentation time of 96 hours. The results are shown in Table 1.

The results shown in Table 1 indicate the importance of aeration in the metabolism of the fungus to provide increased production of the biopolymer. The results using the air flow of 2 vvm after 24 h of process (Experiment 3: 72 h with continuous aeration) showed better results for biopolymer production (14.15 g.L–1), productivity (0.147 g.L–1 h–1), product yield constant (0.89 g of biopolymer/g of glycerol consumed) and glycerol removal (68.34%), among the evaluated experiments.

It can be seen from Table 1 that the aeration promoted in Experiment 3 caused an increase in the production of the biopolymer of 372% compared to Experiment 1 (without addition oxygen), and 33.0% compared to Experiment 2 (with continuous addition of oxygen for 96 h). This shows that there is no need for aeration in the phase where there is no production of gum (first 24 hours of the process). In addition, aeration promotes an increase in the cost of the process. Thus, it was determined that fermentation would be carried out with the addition of air after the 24 hour process.

The results of determination of the effect of yeast extract addition on gum production are shown in Figure 4.

From Figure 4, it is possible to conclude that yeast extract is used by the microorganism as a carbon source to produce the gum, since the increase of the concentration of this compound led to increases in the gum production. Without added yeast extract, the microorganism produced

0.25 gL–1 of biopolymer and it exhibited low consumption of glycerol—only 15.53%.

Souza et al.[19] studied the effect of yeast extract on the biopolymer production by probiotic lactic acid strains, Lactobacillus acidophilus (La-5) and Lactobacillus casei (LC-1). The optimum concentration of the yeast extract obtained was of 0.58% (w/v).

Faria et al.[10] also studied the effect of the addition of yeast extract on the production of xanthana gum by Xanthomonas campestris pv. using broth cane sugar as a carbon source, and found that the optimum condition of the yeast extract concentration was 1.8 g.L–1. As a result, in the next tests, it was chosen the amount of 2 g. L–1 yeast extract used in previous trials, which generated an amount of 14.2 g.L–1 gum and productivity of 0.148 g.L–1 h–1 (Figure 4).

The experiment to evaluate the amount of gum produced using only yeast extract as a carbon source in the medium provided the following results of 0.05 0.15, 0.55 and 1.35 g.L–1 of gum, for the concentration of YE of 1, 2, 3 and 4 g. L–1, respectively. This test confirmed that the fungi applied yeast extract as a carbon source to produce the biopolymer.

3.2 Biopolymer production in fermenter

Previous studies were important to the success of the process in the production of gum in a fermenter. According to Garcia-Ochoa et al.[20], operating conditions such as the configuration of the bioreactor operation mode (batch or

Table 1. Results obtained of formed biopolymer concentration, productivity, product yield constant YP/S and glycerol removal to the performed experiments.

ExperimentsBiopolymer concentration

(g.L–1)Productivity

(g.L–1 h–1)Product yield constant

YP/S

Experiment 1: without additional aeration (96 hours) 3 0.031 0.100Experiment 2: with continuous aeration at 2.0 vvm for 96 hours

10.63 0.111 0.494

Experiment 3: without aeration 24 h and 72 h of process with continuous aeration of 2 vvm

14.15 0.147 0.890

YP/S: g of biopolymer formed/g of glycerol consumed.

Figure 4. Effect of the amount of yeast extract in the formation of the polymer, amount of formed gum (■) and glycerol consumption (○).

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Polímeros, 26(2), 144-151, 2016 149

continuous), medium composition, temperature, pH, agitation speed, aeration rate and the fermentation time influence not only the growth of the microorganism and the production of xanthan gum, but also the structure and rheological properties of the polymer. Thus, the optimization of fermentation conditions on a bench scale is absolutely necessary before the expansion of the production scale.

Once the conditions were set, yeast extract concentration of 2 g.L–1, glycerol concentration of 50 g.L–1, fermentation time of 96 h at 100 rpm, culture time of inoculum of 120 h, and continuous aeration of 2 vvm after 24 h of process, a kinetic study was performed in Biostat B fermenter. Figure 5 shows the curves obtained for cell growth, the amount of formed gum and glycerol consumption as functions of fermentation time.

According to Figure 5, after 16 h of fermentation, the production of biopolymer was 1.1 g.L–1 and productivity of 0.069 g.L–1h–1. Moreover, there is a relationship between cell growth and the amount of gum formed, since the amount of gum formed increased with the growth of the microorganism.

In 74 h there was a maximum production of gum, 16.35 g.L–1 (productivity of 0.221 g.L–1h–1) with a maximum cell growth of 6.5 g.L–1. Another finding was the reduction in the content of biopolymer after 74 h, along with an increase of glycerol consumption. After 74 h of fermentation, it was verified visually that the fibers were shattered, presenting a different structure from those fibers produced in 74 h of process where fibers were compact and long. This fact indicates that the fungus

could have produced enzymes capable of degrading the gum using the remaining glycerol as a substrate for production of other by-products.

Faria et al.[10], studied the kinetic profile of gum production using Xanthomonas campestris pv. campestris NRRL B-1459 in the bioreactor and it obtained 16.4 g.L-1 of gum.

The values found in this study were next compared to Faria et al.[10], as was shown above, proving that the study of the variables is extremely important to the produced biopolymer.

Knowing about the properties of polymers, especially the viscosity and rheological behavior, are important for future industrial applications because it allows us to obtain information about the deformation and flow properties of materials[21].

The rheological properties were evaluated by the apparent viscosity analysis to verify the quality of the gum produced in the previously selected conditions. Figure 6 shows the behavior of viscosity of the gum as a function of shear rate for different gum concentrations (0.5, 0.75 and 1%). According to Faria et al.[12], an important feature of biopolymers is their ability to cause thickening at a given concentration, generally less than 1% (w/v), promoting higher viscosities to the known strain rates.

The results indicated that the apparent viscosity decreased with increasing the shear rate. According to the literature, this behavior has been found in polymer solution microbial polysaccharides[22].

Beyer et al.[23] found that 1% solutions of polysaccharide produced by Rhizobium CB744 at 25 °C, also promoted a decrease in viscosity with increasing shear rate.

In the comparative graphs of the apparent viscosity readings of aqueous solutions of gum produced, it can be seen that the concentration of 0.5% has the lowest viscosity values, of 143 mPa.s at a shear rate of 4 s–1 (Figure 6).

The solutions of concentrations from 0.75% to 1% at (under) shear rate of 4 s–1, had viscosities of about 2,730 to 1,010 mPa.s, respectively (Figure 6). This shows that at a low shear rate the produced gum showed higher viscosity.

3.2.1 Characterization of the formed goma: determination of monosaccharide composition and molecular parameters

The presence of amino sugars was investigated by thin layer chromatography (TLC). The patterns of amino sugars (glucosamine, and mannosamine) were revealed with ninhydrin and orcinol. This was negative for the presence of amino sugars in the sample.

Table 2 shows the monosaccharide composition of the gum formed by the fungus Mucor racemosus Fresenius.

Figure 5. Profile Kinetic of the biopolymer production in the bioreactor, amount of formed gum (●), glycerol consumption (○) and cell growth (♦).

Figure 6. Variation of viscosity versus shear rate of aqueous solutions of the polymer in concentrations of 0.5; 0.75 and 1% at 25 °C.

Table 2. Monosaccharide composition of the formed gum.Components Composition (%)

Rhamnose 6.6 ± 0.52Manose 28.3 ± 1.79Galactose 32.1 ± 0.81Glucose 22.5 ± 1.08Glucuronic acid 10.0 ± 0.01Galacturonic acid 0.5 ± 0.01

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As Table 2 shows, the total sugar content for the biopolymer produced was 89.5%. It can be seen that the gum is composed of larger amounts of sugars mannose, galactose and glucose, and minor amounts of rhamnose and uronic acids.

TLC confirmed the presence of neutral sugars and indicated that the main sugar acid is glucuronic acid at 10%. The ratio of components glucose, mannose and glucuronic acid obtained in the present study was 2.25: 2.83: 1. The ratio of these components (glucose, mannose and glucuronic acid) in the xanthan gum obtained by Rosalam and England[24] was 2: 2: 1 and Faria et al.[12] was 1.79: 1:33: 1. The ratios different found in this paper. Furthermore, the gum obtained in this paper presents sugars (rhamone and galactose) that are not in significant quantities in the structure of the xanthana gum. Therefore, the gum produced in the present study using of the cerrado fungus (Mucor racemosus Fresenius) and glycerol as raw material, showed different structure from the Xanthan gum.

Analyzing the results of unimodal distribution of mass obtained was observed value of polydispersity index (Mw / Mn) equal to 1.002 (Table 3). This polydispersity value shows homogeneity in the analyzed material.

The molecular parameters obtained to the sample from the elution profile using the Zimm method are shown in Table 3.

The weighted average molecular weight (Mw) obtained of the gum was 4.607×106 g.mol–1 (Da). By analyzing this parameter, it can be seen that this value is within the molecular weight range provided for the xanthan gum can vary from 5.0×107 to 2×106 Da (Daltons)[25].

The variations in the fermentation conditions are factors that influence the molar mass of xanthan. The use of xanthan gum as a reference if should the interest of the physico-chemical properties that surpass all other polysaccharides available. The property that stands out is its high viscosity at low concentrations (0.05-1%), being induced by their branched structure and high molecular weight and stability in a wide range of temperature and pH[20,26].

4. Conclusions

Glycerol showed a promising substrate for the production of biopolymer using the fungus Mucor racemosus Fresenius. The results showed production of 14.15 g.L–1 gum and consumption of 68.34% glycerol at selected conditions –50 g.L–1 glycerol in the initial table shaker at 100 rpm, with air

flow of 2 vvm during production (without aeration in the first 24 hours and with continuous aeration in 72 hours of assay), age inoculum of 120 hours and room temperature of 28 ± 2 °C. The use of aeration increased the production of gum to 371%. Under conditions previously determined, assays carried out in a reactor of 2 L with a useful volume of 1 L, at 74 h of processing obtained 16.35g.L–1 gum (productivity of 0.221 g.L–1 h–1) and 75% glycerol consumption. The gum presented a pseudoplastic behavior with higher viscosity than the xanthan gum in the shear rate of 4 s–1, similarly to the biopolymer reported by the literature. The weighted average molecular weight (Mw) obtained for the gum was 4.607 × 106 Da. The value of the polydispersity index (Mw/Mn) equal to 1.002 shows the homogeneity of the produced biopolymer. The monosaccharide composition of the polymer was 6.6% of rhamnose, mannose 28.3%, 32.1% galactose, 22.5% glucose, 10% glucuronic acid and 0.5% galacturonic acid. It is worth mentioning that the production of a large amount of biopolymer using the fungus Mucor racemosus Fresenius and employing glycerol as substrate is not easily found in the literature.

5. Acknowledgements

This research was funded by Fundação de Amparo a pesquisa do Estado de Minas Gerais (FAPEMIG), Conselho Nacional de Desenvolvimento Científico e Tecnológico (CNPQ) and Coordenação de Aperfeiçoamento de Pessoal de Nível Superior (CAPES).

6. References

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6. Alves, M. H., Takaki, G. M. C., Okada, K., Pessoa, I. H. F., & Milanez, A. I. (2005). Detection of extracellular protease in Mucor species. Revista Iberoamericana de Micologia, 22(2), 114-117. http://dx.doi.org/10.1016/S1130-1406(05)70020-6. PMid:16107171.

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Table 3. Molecular parameters of the formed gum.Variables Results

Mw 4.607 × 106 g.mol–1

Mn 4.600 × 106 g.mol–1

Mz 4.613 × 106 g.mol–1

Polydispersity index (Mw/ Mn) 1.002 ± 0.011

Rn 56.0 nm

Rw 56.3 nm (error 1.7%)

Rz 56.5 nm (error 1.7%)

Biopolymer production using fungus Mucor racemosus Fresenius and glycerol as substrate

Polímeros, 26(2), 144-151, 2016 151

racemosus. Bioresource Technology, 97(3), 506-511. http://dx.doi.org/10.1016/j.biortech.2005.02.046. PMid:15979307.

8. Amorim, R. V. S., Souza, W., Fukushima, K. E., & Campos-Takaki, G. M. (2001). Faster chitosan production by mucoralean strains in submerged culture. Brazilian Journal of Microbiology, 32(1), 20-33.

9. Faria, S., Vieira, P. A., de Resende, M. M., Ribeiro, E. J., & Cardoso, V. L. (2010). Application of a model using the phenomenological approach for prediction of growth and xanthan gum production with sugar cane broth in a batch process. LWT – Food Science and Technology (Campinas.), 43, 498-506. http://dx.doi.org/10.1016/j.lwt.2009.09.018.

10. Faria, S., Vieira, P. A., Resende, M. M., França, F. P., & Cardoso, V. L. (2009). A comparison between shaker and bioreactor performance based on the kinetic parameters of xanthan gum production. Applied Biochemistry and Biotechnology, 156(1-3), 475-488. http://dx.doi.org/10.1007/s12010-008-8485-8. PMid:19130306.

11. Freitas, F., Alves, V. D., Pais, J., Costa, N., Oliveira, C., Mafra, L., Hilliou, L., Oliveira, R., & Reis, M. A. (2009). Characterization of an extracellular polysaccharide produced by a Pseudomonas strain grown on glycerol. Bioresource Technology, 100(2), 859-865. http://dx.doi.org/10.1016/j.biortech.2008.07.002. PMid:18713662.

12. Faria, S., Petkowicz, C. L. O., Morais, S. A. L., Terrones, M. G. H., Resende, M. M., França, F. P., & Cardoso, V. L. (2011). Characterization of xanthan gum produced from sugar cane broth. Carbohydrate Polymers, 86(2), 469-476. http://dx.doi.org/10.1016/j.carbpol.2011.04.063.

13. Biermann, C. J., & Mcginnis, G. D. (1989). Analysis of carbohydrates by GLC and MS. Florida: CRC Press.

14. Blumenkrantz, N., & Asboe-Hansen, G. (1973). New method for quantitative determination of uronic acids. Analytical Biochemistry, 54(2), 484-489. http://dx.doi.org/10.1016/0003-2697(73)90377-1. PMid:4269305.

15. Nagel, A., Sirisakulwat, S., Carle, R., & Neidhart, S. (2014). An acetate−hydroxide gradient for the quantitation of the neutral sugar and uronic acid profile of pectins by HPAEC-PAD without postcolumn pH adjustment. Journal of Agricultural and Food Chemistry, 62(9), 2037-2048. http://dx.doi.org/10.1021/jf404626d. PMid:24547908.

16. Chaplin, M. F. (1994). Carbohydrate analysis: a practical approach. London: IRL Press Practical Approach Series.

17. Antônio, R. V., Recouvreux, D. O. S., Nazario, A. C., Timboni, D., Ferrarini, E., Rodowanski, G. P., Cauduro, M. T., & Peres, S. S. (2012). Produção de celulose bacteriana a partir de diferentes substratos. Revista Técnico Científica, 3(1), 176-182.

Retrieved in 19 August 2015, from http://periodicos.ifsc.edu.br/index.php/rtc/article/view/726

18. Souza, A. S., & Vendruscolo, C. T. (2000). Produção e caracterização dos biopolímeros sintetizados por Xanthomonas campestris pv. Pruni cepas 24 e 58. Ciência e Engenharia, 8, 115-123.

19. Souza, T. D. S., Yuhara, T. T., Castro-Gómez, R. J. H., & Garcia, S. (2007). Produção de exopolissacarídeos por bactérias probióticas: otimização do meio de cultura. Brazilian Journal of Food Technology, 10(1), 27-34. Retrieved in 19 August 2015, from http://www.ital.sp.gov.br/bj/artigos/bjft/2007/p06269.pdf

20. Garcia-Ochoa, F., Castro, E. G., & Santos, V. E. (2000). Oxygen transfer and uptake rates during xanthan gum production. Enzyme and Microbial Technology, 27(9), 680-690. http://dx.doi.org/10.1016/S0141-0229(00)00272-6. PMid:11064050.

21. Scamparini, A., Mariuzzo, D., Fujihara, H., Jacobusi, R., & Vendruscolo, C. (1997). Structural Studies of CV-70 Polysaccharide. International Journal of Biological Macromolecules, 21(1-2), 115-121. http://dx.doi.org/10.1016/S0141-8130(97)00050-0. PMid:9283025.

22. Amanullah, A., Serrano, L. C., Galindo, E., & Nienow, A. W. (1996). Reproducibility of pilot scale xanthan fermentations. Biotechnology Progress, 12(4), 466-473. http://dx.doi.org/10.1021/bp960042k.

23. Beyer, R., Melton, D. L., & Kennedy, D. L. (1987). Viscosity studies on the polysaccharide gum from Rhizobium strain CB 744. Journal of the Science of Food and Agriculture, 39(2), 151-161. http://dx.doi.org/10.1002/jsfa.2740390208.

24. Rosalam, S., & England, R. (2006). Review of xanthan gum production from unmodified starches by Xanthomonas camprestris sp. Enzyme and Microbial Technology, 39(2), 197-207. http://dx.doi.org/10.1016/j.enzmictec.2005.10.019.

25. Papagianni, M., Psomas, S. K., Batsilas, L., Paras, S. V., Kyriakidis, D. A., & Liakopoulou-Kyriakides, M. (2001). Xanthan production by Xanthomonas campestris in batch cultures. Process Biochemistry, 37(1), 73-80. http://dx.doi.org/10.1016/S0032-9592(01)00174-1.

26. Diaz, O. S., Vendruscolo, C. T., & Vendruscolo, J. L. S. (2004). Reologia de Xantana: uma revisão sobre a influência de eletrólitos na viscosidade de soluções aquosas de gomas xantana. Semina: Ciências Exatas e Tecnológicas, 25(1), 15-28. Retrieved in 19 August 2015, from http://www.uel.br/revistas/uel/index.php/semexatas/article/viewFile/1557/1308

Received: Aug. 19, 2015 Revised: Nov. 06, 2015

Accepted: Nov. 09, 2015

http://dx.doi.org/10.1590/0104-1428.2053

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Influence of nucleating agent on the crystallization kinetics and morphology of polypropylene

Adriane Gomes Simanke1*, Ana Paula de Azeredo1, Cristóvão de Lemos1 and Raquel Santos Mauler2

1Braskem S.A., Triunfo, RS, Brazil2Instituto de Química, Universidade Federal do Rio Grande do Sul – UFRGS, Porto Alegre, RS, Brazil

*adriane.simanke@braskem.com

Sbstract

The influence of three nucleating agents from different generations on the crystallization behavior of propylene homopolymer was studied by differential scanning calorimetry (DSC) and atomic force microscopy (AFM). The amount of nucleating agent used varied between 1000 and 2200 ppm. The new generation nucleating agent, Hyperform HPN-68L, accelerates the crystallization more efficiently than the other nucleating agents tested. It was also possible to verify the effects of agglomeration and negative interaction between calcium stearate and sodium benzoate. Furthermore, AFM images allowed to differentiate the crystals generated by Millad 3988 through the observation of a fibrillar intertwining network structure, with characteristic spacing and length of crystals, justifying its excellent performance to improve polypropylene optical properties.

Keywords: AFM, crystallization kinetics, morphology, nucleating agent, polypropylene.

1. Introduction

Propylene resins are one of the most versatile families of thermoplastics. Their semicrystalline nature and their structural stereoregularity allow controlling their physical and mechanical properties through a variety of variables such as molecular structure, molecular weight and molecular weight distribution, comonomer content and distribution, degree of crystallinity and morphology. Advances in the catalytic system, polymerization process and post reactor changes, including the use of different additives, allow to overcome polypropylene previous limitations, making it competitive with other polymers and materials. Polypropylene and its copolymers are widely used in applications ranging from durable parts such as exterior and interior parts of automobiles to disposable packaging such as bags. Although polypropylene faces some restrictions in applications that require outstanding optical properties, these can be improved by propene copolymerization with comonomers such as ethylene and butene or by the use of additives, as the nucleating agents. These foreign particles (as silica, talc and organic salts) added to polypropylene act as seeds of nucleation process, accelerating the crystallization rate[1]. Nucleating agents are being widely used in order to improve mechanical and optical properties of polypropylene and its copolymers. They are also used to accelerate polypropylene crystallization kinetics, reducing injection molding cycle times and, by consequence, reducing production costs. The nucleating agent efficiency depends on its particle size, morphology, chemical structure and behavior when incorporated into the polymer[2]. In order to maximize the nucleation efficiency and reach the best cost benefit ratio, nucleating agents evolved and there are various types in the market, from different chemical families. Although a large number of papers have been published regarding the use of nucleating agents in polypropylene, just a few[3-5] describe in detail the influence of different nucleating agent structures

on polypropylene morphology and crystallization kinetics. In this work, three commercial nucleating agents from different generations were evaluated and their influence on the crystallization kinetics and morphology were studied through atomic force microscopy (AFM) and differential scanning calorimetry (DSC).

2. Materials and Methods

2.1 Materials

A propylene homopolymer (PP) produced by Braskem (see characteristics in Table 1) was used to evaluate the performance of three commercial nucleating agents: 1,3:2,4-bis(3,4-dimethylbenzylidene)sorbitol (Millad 3988 from Milliken), identified as NA, sodium benzoate, identified as NB, and sodium bicycle[2,2,1]heptane dicarboxylate salt (Hyperform HPN-68L from Milliken), identified as NC. At least three different concentrations of these nucleating agents were used: 1000, 1600 and 2200 ppm (Table 2).

2.2 Sample preparation

Polypropylene powder and nucleating agents were mixed in an intensive mixer, Mixaco CM-600D, at 1600 rpm, under nitrogen for 30 s. It was also added 1200 ppm of Irganox B-215 and 500 ppm of calcium stearate (CaSt). Then, samples were extruded in a single screw extruder, Rulli EF-70, L/D 1:25, with 70mm screw diameter. Temperatures of the five zones of the extruder were 180, 185, 190, 195, 200 °C and the extruder screw speed was adjusted to 130 rpm.

Samples were compressed molded at 175 °C/30 ton in a hydraulic press (G302 Wabash equipment) to obtain films of 35-45 µm thickness to be used in the DSC and AFM analyses.

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2.3 Characterization techniques

2.3.1 Differential Scanning Calorimetry (DSC)

Thermal behavior was analyzed by differential scanning calorimetry using a TA Q1000 DSC under nitrogen and connected to an intracooler that allowed sub ambient temperature control. The instrument was calibrated with indium. Film samples (6 ± 1 mg) were melted at 200 °C, held at this temperature for 5 minutes to ensure complete melting, cooled from 200 °C to -20 °C and heated from –20 °C to 200 °C. All heating/cooling rates used were 10 °C/min. The melting temperature (Tm) and enthalpy of fusion (ΔHf) were taken from the second heating curve. Crystallinity degree (Xc) is given byΔHf / H

0f , where H0

f is the enthalpy of fusion of 100% crystalline PP (190 J/g).

The isothermal crystallization experiments were carried out in a TA Q1000 DSC equipped with a liquid nitrogen cooling system. Samples were heated to 200 ºC, held at this temperature for 5 minutes to ensure complete melting and quenched to the selected isothermal crystallization temperature at a nominal cooling rate of 140 ºC/min. In order to calculate the Avrami coefficients, three different crystallization temperatures were used: 146.5 °C, 147 °C and 147.5 °C.

Analyses of the experimental time-dependent relative crystallinity function θ(t) were carried out according to the Avrami model[6-9], which is given by:

θ(t) = 1 - exp(-ktn) (1)

where k is the Avrami rate constant, and n is the Avrami exponent. Both k and n are constants specific to a given crystalline morphology and type of nucleation for a particular crystallization condition[7]. Another parameter to evaluate the nucleating performance is the crystallization half time (t1/2)

[8], which is a measure of the time it takes from the onset of crystallization until the crystallization is half completed.

2.3.2 Scanning Electron Microscopy (SEM)

The morphology of the nucleating agent surface was investigated with scanning electron microscopy (Hitachi TM-1000 model). Samples were analyzed in the powder form without previous sputter-coated and the electron micrographs were taken using an acceleration voltage of 15.0 kV.

2.3.3 Atomic Force Microscopy (AFM)

AFM images were obtained using Veeco NanoScope V atomic force microscopy and Veeco diTAC heater system operating under heating conditions. Topography and phase images were simultaneously collected in tapping mode at 512 x 512 lines standard resolution. Veeco single side coated silicon cantilevers were used with resonant frequency at 366-401 kHz. According to the manufacturer’s specifications, the cantilevers have a spring constant of 20-80 N/m, length of 110-140 µm, width of 25-35 µm and the radius of the tip is 2-5 nm. Film samples of nominal thickness of 35-45 µm and without previous surface treatment were subject to different scan areas using a scan rate between 0.0313 to 0.878 Hz. Films were quickly heated to 200 °C and kept at this temperature for 5 min. In a second stage, the melted polymer was quickly cooled to the isotherm temperature of scanning to observe the crystallization step and lamellae growth. PP NA 1, PP NB 1 and PP NC 1 were analyzed at 155 °C whereas neat PP was analyzed at 145 °C and the isothermal crystallization time varied according to the crystallization rate.

2.3.4 Optical properties - Haze and gloss 45°

Haze measurements were carried out on a Haze Gard Plus made by BYK-Gardner, according to ASTM D 1003-00. Gloss measurements were carried out on a Micro Gloss 45° made by BYK-Gardner, using light source C, according to ASTM D 2457-03. Haze and gloss were measured using 1 mm x 60 mm x 60 mm injection molded plates. Injection molded plates were obtained using an Arburg Allrounder 270U 400-170 equipment at 170/175/180/185/190°C barrel temperatures, 300 cm3/s injection flow, 60°C mold temperature (procedure according to ASTM D-3641). After injection molding, specimens were condicioned at 23 ± 2 ºC and 50 ± 5% RH for 40 h.

3. Results and Discussions

3.1 Nucleating agent morphology

Scanning Electron Microscopy (SEM) analyses of the three nucleating agents used in this work were carried out in order to evaluate their morphology (Figure 1).

The nucleating agent NA shows fibrillar morphology whereas in NB and NC the predominant morphology is granular. Moreover, the granules of NC are smaller and more homogeneous than NB. It is also possible to observe the presence of some agglomerated particles in NB due to its hygroscopicity[10].

3.2 Thermal analysis characterization

Comparing neat polypropylene (PP) with nucleated PP samples (Table 3), it is possible to observe that the addition of 1000 ppm of nucleating agent is sufficient to cause a significant increase in the crystallization temperature (Tc).

Table 2. Samples composition.PP neat PP

PP NA 1 PP + 1000 ppm Millad 3988PP NA 2 PP + 1600 ppm Millad 3988PP NA 3 PP + 2200 ppm Millad 3988PP NB 1 PP + 1000 ppm sodium benzoatePP NB 2 PP + 1600 ppm sodium benzoatePP NB 3 PP + 2200 ppm sodium benzoatePP NC 1 PP + 1000 ppm Hyperform HPN-68LPP NC 2 PP + 1600 ppm Hyperform HPN-68LPP NC 3 PP + 2200 ppm Hyperform HPN-68L

Table 1. Polypropylene characteristics.

SampleMFI

230°C/2.16kg Density Mn MwMw/Mn

(g/10 min) (g/cm3) (kg/mol) (kg/mol)PP 20 0.905 39200 163700 4.2

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Among the three nucleating agents tested, the highest increase in Tc is obtained with NC. It is noted, however, that the increase of Tc is not proportional to the increment in the amount of nucleating agent. Crystallization temperatures of the samples with 1600 ppm and 2200 ppm of NA and NC are only 1 °C or 2 °C higher than the Tc of the samples with 1000 ppm, indicating that the lowest amount tested could be enough to improve the thermal properties. In the samples with NB, it is observed a decrease in Tc as the nucleating amount increases. This behavior indicates that the amount of NB used is above the ideal concentration to attain the highest nucleation rate and the excess of additive could be causing agglomeration, decreasing nuclei number and, by consequence, decreasing Tc[11]. Libster et al.[12] showed that a good dispersion of NC in microemulsions causes super saturation since low amounts, indicating that higher amounts do not supply additional nuclei, that could explain the similar Tc values obtained for the different amounts of NC tested.

In order to verify if nucleating agent levels below 1000 ppm also increase the crystallization temperature efficiently, samples with lower amounts of NC and NB were prepared and analyzed. Figure 2 shows the crystallization temperature as a function of the nucleating agent amount for PP samples nucleated with NC and NB.

Analyzing Figure 2, it is possible to see that there is an increase of Tc with increasing content of NC. Although the increase is small, it cannot be attributed to a poor dispersion, but to the fact that the maximum performance of NC is achieved at low concentrations[12]. In the case of NB, it is possible to observe that the Tc increases as the nucleating agent amount increases from 200 ppm to 1000 ppm, but above this concentration, it decreases. This behavior is in accordance with the agglomeration hypothesis that has been already discussed by Botkin et al.[13].

3.2.1 Isothermal crystallization kinetics

The nucleated samples were isothermally crystallized at three different temperatures: 146.5 °C, 147 °C and 147.5 °C. Because neat PP did not crystallize at these temperatures, it was analyzed at 134 °C and 136 °C. The rate of evolution of the crystallization enthalpy with time was measured by DSC and the integration of dHc/dt in the crystallization range is expressed as the relative crystallinity (Xt) as a function of time.

Since each nucleated sample showed similar behavior at the three different crystallization temperatures, only the crystallization kinetics curves obtained at 146.5 °C for each nucleated sample will be presented.

Figure 3 shows the relative crystallinity curves as a function of time at 146.5 °C for the three PP samples nucleated with NA. As the amount of NA is increased, it is observed a substantial increase in the crystallization rates, which is in accordance with some studies using sorbitol derivatives, such as NA, that indicate that they need a

Figure 1. SEM images of nucleating agents (a) NA, (b) NB and (c) NC.

Table 3. Thermal Properties of Neat PP and Nucleated PP Samples.Samples Tc (°C) Tm2 (°C) Xc2 (%)

PP 113 162 58PP NA 1 130 165 56PP NA 2 132 165 58PP NA 3 132 165 60PP NB 1 132 165 59PP NB 2 130 165 57PP NB 3 128 164 55PP NC 1 135 165 59PP NC 2 136 165 58PP NC 3 136 165 60

Figure 2. Crystallization Temperature as a Function of the Amount of Nucleating Agent.

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Polímeros, 26(2), 152-160, 2016 155

minimum concentration to begin having nucleating effects on polypropylene[14-16].

Regarding PP samples nucleated with NB, it is possible to see that the crystallization rate decreases as the amount of NB agent is increased, which can be explained by the agglomeration of the additive that decreases the number of nuclei formed. Similar behavior is observed at all temperatures analyzed (Figure 4).

In the case of samples nucleated with NC, the increase in the crystallization rate is not very significant as the amount of NC is increased, indicating that its maximum performance in terms of thermal properties is achieved with low concentrations (Figure 5).

Comparing the three nucleating agents, it is possible to conclude that the highest crystallization rate is obtained with NC, in the three contents tested (1000 ppm, 1600 ppm and 2200 ppm). NC accelerates the crystallization more efficiently than the other nucleating agents tested. The crystallization rates of samples nucleated with NB are intermediary to the ones achieved by the samples nucleated with NC and NA.

3.2.2 Avrami parameters

The Avrami coefficients of neat PP were calculated in only two temperatures and the n values obtained (Table 4) were in agreement with values found in the literature[6,9] for polypropylene. The Avrami coefficients of the nucleated samples were calculated at the three temperatures studied and the results are shown in Table 5. The n values obtained for the nucleated samples are around 3.0, ranging between 2.1 and 3.4, pointing to the formation of crystals varying between two-dimensional and three-dimensional structures. It was expected the decrease of k values with the increase of temperature, but some of the samples did not follow this tendency.

Regarding the crystallization half time (t1/2), it is possible to observe the increase of t1/2 values with the increase of NB content, confirming the agglomeration problem. It is observed a significant decrease in t1/2 values as the amount of NA was increased from 1000 ppm to 2200 ppm, while t1/2 values are similar among the different levels of NC used.

Based on these values, one can conclude that NC accelerates the crystallization rate more effectively than the other nucleating agents studied.

3.3 AFM characterization

AFM characterization was carried out at different temperatures in order to define the better isothermal temperature to analyze the specimens. AFM height and phase images of PP obtained at 2 and 17 minutes at an isotherm of 145 °C are showed in Figure 6.

In Figure 6a it is observed some dots indicating the presence of crystalline nuclei that characterizes the

crystallization beginning. Then, molecules start to reorganize themselves and they change from an intermediate degree of disorder to a more organized crystalline state. In the sequence, crystals continue to grow and after 17 minutes of isotherm it is verified the presence of regular and well defined spherulites with equivalent diameter (18 ± 1 µm) and similar morphology. Although neat PP shows crystals with homogeneous morphological characteristics, they are not very well distributed since there are regions without crystals and others where crystals are overlapped.

AFM characterization of nucleated samples was carried out at 155 °C. Figure 7 shows crystal growth and morphology over the time for PP NB 1. As expected, spherulite size is reduced compared to neat PP. The images allow visualizing the spherulite formation architecture, being

Figure 3. Isothermal crystallization curves of NA nucleated samples, at 146.5 °C.

Figure 4. Isothermal crystallization curves of NB nucleated samples, at 146.5 °C.

Figure 5. Isothermal crystallization curves of NC nucleated samples, at 146.5 °C.

Table 4. Avrami parameters for neat PP at different isothermal temperatures.

Parameters n K t1/2 (min)

Temperature (°C)

134 136 134 136 134 1362.5 2.8 9.23 e-05 6.11 e-05 30.1 32.9

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able to distinguish the first crystals formed at 7 minutes. At 11 minutes of isotherm, it is observed the presence of similar size lamellar sheaves randomly distributed that grow oriented away from the nucleus, followed by slower in-filling ‘secondary’ lamellar growth[17]. The presence of a contrast light gray indicates a melt region, whereas the dark gray represents the crystalline region.

Figure 8 shows the morphology and kinetic profile for PP NC 1 isothermally crystallized at 155 °C. After 4 minutes of isotherm, it can be observed the presence of a large amount of embryos/nuclei satisfactory dispersed. At the height images it is clearly observed some lighter regions indicating that the crystals emerge from the molten phase, grow and form branches from the initial crystals. After 13 minutes

Figure 6. AFM images of PP isothermally crystallized at 145 °C (a) 2 min. and (b) 17 min.

Table 5. Avrami parameters for nucleated PP at different isothermal temperatures.

Parameters n Kt 1/2 (min)

(min)Temperature 146.5 °C 147 °C 147.5 °C 146.5 °C 147 °C 147.5 °C 146.5 °C 147 °C 147.5 °C

PP NA 1 3.1 3.4 2.7 2.51E-06 7.94E-07 3.98E-05 47.4 54.7 53.7PP NA 2 3.0 2.7 3.2 2.51E-05 1.00E-04 5.01E-06 28.8 27.3 40.7PP NA 3 2.1 2.1 2.2 2.51E-03 1.58E-03 6.31E-04 16.7 19.1 24.6PP NB 1 3.1 3.0 3.4 6.31E-04 6.31E-04 1.58E-05 9.5 10.2 11.8PP NB 2 3.1 3.0 3.1 1.26E-04 1.58E-03 2.51E-05 14.9 16.3 20.0PP NB 3 3.5 3.2 3.4 1.58E-05 3.16E-05 1.58E-05 21.3 22.0 24.0PP NC 1 3.0 3.2 3.1 1.00E-02 6.31E-03 3.16E-03 4.3 4.8 5.9PP NC 2 2.9 2.7 2.9 1.58E-02 2.00E-02 1.00E-02 3.7 3.8 4.4PP NC 3 2.7 2.9 3.1 2.51E-02 1.58E-02 7.94E-03 3.4 3.8 4.4

Figure 7. AFM images of PP NB 1 isothermally crystallized at 155 °C (a) 7 min., (b) 11 min. and (c) 22 min.

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Polímeros, 26(2), 152-160, 2016 157

of isotherm, more than 50% of the sample is crystallized. NC produces crystals with smaller size and faster kinetic profile when compared to the other nucleating agents. These observations are in agreement with DSC crystallization kinetics studies, confirming the faster crystallization rate showed by the samples nucleated with NC.

Comparing AFM images of PP NB 1 and PP NC 1 obtained at 155 °C after 22 minutes and 13 minutes of isotherm, respectively, (Figure 9a and Figure 9b) it is possible to observe that both show similar morphology. However, the presence of NC contributes to reduce the size of the crystals and increase the number of crystallization nuclei.

Figure 10 shows the AFM images of sample PP NA 1, where it is not observed the characteristic spherulite

morphology or other similar structures. It is verified the presence of a very thin and unique fibrillar structure intertwining networks without the formation of specific nuclei. Similar structures were also observed in some studies carried out with derivatives of sorbitol nucleating agents[15,17,18]. Probably this morphology is responsible for the improvement in the optical properties, reducing the interaction between the polymer crystalline phase and visible light. Due to this behavior, NA can also be applied as clarifier.

3.4 Optical properties

As expected, nucleated samples show optical properties improvements. The exceptions are the samples nucleated with NB, as can be observed in Figure 11 and 12. Samples

Figure 8. AFM images of PP NC 1 isothermally crystallized at 155 °C (a) 4 min., (b) 9 min. and (c) 13 min.

Figure 9. Comparison between spherulite size of sample (a) PP NB 1 (22 minutes of isotherm, 155 °C) and (b) PP NC 1 (13 minutes of isotherm, 155 °C).

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nucleated with NA show the best improvement in optical properties, as expected for a clarifier. This better performance can be explained by its fibrillar structure, as could be seen in AFM images (Figure 10). Clarifying agents such as NA dissolve and mix with molten PP during melt processing. Upon cooling, the clarifying agent crystallizes first in the form of a three-dimensional nanometric fibrillar network and acts as a clarifying agent for PP. This fibrillar network ensures extremely fine dispersion of the nuclei and can facilitate the subsequent processes of nucleation and crystal growth[12]. The crystallites obtained are so small that incident light scattering is significantly decreased and, by consequence, optical properties are improved. As the amount of NA increases, gloss improves gradually and haze decreases. A maximum of 40% increase of gloss and a maximum of 70% reduction in haze are achieved compared to neat PP. It is observed an increase of 17% in gloss and a decrease of 36% in haze with NC addition, being the improvements almost the same for the three amounts tested. When NB is added, it is observed a decrease in optical properties, about 3% in gloss and an increase of 20% in haze. Although samples nucleated with NB show a reduction in crystal size compared to neat PP (Figures 6 and 7), there is an increase of haze. This negative effect on optical properties could be consider as a consequence of the agglomeration caused by the excess of NB, but even samples nucleated with 1000 ppm of NB (the optimum amount tested for NA) did not show an improvement in these properties.

PP nucleation studies[19,20] show that calcium stearate, used as anti-acid in these formulations, may interfere in NB action, decreasing its performance as nucleating agent. Some studies[10,19] propose an ion interchange between calcium and sodium, forming compounds that are not nucleating agents. However, Dieckmann[20] showed that this negative effect occurs not by the calcium presence but by the presence of a metallic ion stearate. It was also demonstrated that the use of sodium stearate in samples nucleated with NB do not prejudice the efficiency of the nucleating agent.

Figure 10. AFM images of PP NA 1 isothermally crystallized at 155 °C (a) 5 min., (b) 9 min. and (c) 11 min.

Figure 11. Gloss results of neat and nucleated PP.

Influence of nucleating agent on the crystallization kinetics and morphology of polypropylene

Polímeros, 26(2), 152-160, 2016 159

In order to check the influence of calcium stearate in the performance of NB, it was produced one sample with 1000 ppm of NB without calcium stearate. Thermal and optical properties were measured and the results are shown in Table 6. Although there are no differences in crystallization temperature, it is observed an improvement in optical properties with the absence of calcium stearate. A reduction of about 15% in haze is obtained when compared to the sample containing calcium stearate, confirming the negative interaction.

As NC is also a sodium salt and do not melt into the polymer mass, like NB, it was prepared a sample with 1000 ppm of NC without calcium stearate in order to verify if there is also an interference between these additives. Analyzing Table 7, it is not verified the same negative influence. The sample without calcium stearate shows a decrease in Tc and gloss and an increase in haze, when compared to the sample with calcium stearate. This result is according to data showed in the patent US2006/0142452 A1[20], where different stearates were tested in formulations with NC and it was observed an optimization of NC performance when used with calcium or zinc stearate.

4. Conclusions

PP nucleated samples showed different characteristics and behavior according to the nucleating agent type and amount used. The old generation and more commom nucleating agent tested (sodium benzoate) showed some limitations regarding agglomeration and interaction with stearates, but it is a good choice to get improvements in PP thermal properties. Hyperform HPN-68L, an advanced nucleating agent, showed to be more efficient to accelerate crystallization, even with low amounts. Also, it was possible to understand why PP nucleated with Millad 3988 shows very good improvements in optical properties through AFM images, where a network of fibrillar structure was observed instead of defined spherulites.

It was also possible to verify that there is an optimum amount of nucleating agent to reach the best performance and properties in nucleated PP and it depends on the type of nucleating agent. In many cases, the use of higher amounts of these additives is unnecessary and, even, as in the case of sodium benzoate, it can affect optical properties.

5. Acknowledgements

Authors would like to thank Braskem S.A. and UFRGS for supporting this work.

6. References

1. Mubarak, Y., Harkin-Jones, E., Martin, P. J., & Ahmad, M. (1999). Crystallization of isotactic polypropylene: pigment, nucleating agent and recycling effects. In Proceedings of 57° SPE ANTEC (pp. 3796-3800). New York: SPE.

2. Wang, K., Mai, K., & Zeng, H. (2000). Isothermal crystallization behavior and melting characteristics of injection sample of nucleated polypropylene. Journal of Applied Polymer Science, 78(14), 2547-2553. http://dx.doi.org/10.1002/1097-4628(20001227)78:14<2547::AID-APP160>3.0.CO;2-F.

3. Menyhárd, A., Gahleitner, M., Varga, J., Bernreitner, K., Jääskeläinen, P., Øys, H., & Pukánszky, B. (2009). The influence of nucleus density on optical properties in nucleated isotactic polypropylene. European Polymer Journal, 45(11), 3138-3148. http://dx.doi.org/10.1016/j.eurpolymj.2009.08.006.

4. Marco, C., Ellis, G., Gomez, M. A., & Arribas, J. M. (2003). Analysis of the isothermal crystallization of isotactic polypropylene nucleated with sorbitol derivatives. Journal of Applied Polymer Science, 88(9), 2261-2274. http://dx.doi.org/10.1002/app.11935.

5. Binsbergen, F. L., & Lange, B. G. M. (1970). Heterogeneous nucleation in the crystallization of polyolefins: Part 2. Kinetics of crystallization of nucleated polypropylene. Polymer, 11(6), 309-332. http://dx.doi.org/10.1016/0032-3861(70)90071-6.

6. Hay, J. N. (1971). Application of the modified avrami equations to polymer crystallisation kinetics. Polymer Journal, 3(2), 74-82. http://dx.doi.org/10.1002/pi.4980030205.

7. Supaphol, P. (2001). Application of the Avrami, Tobin, Malkin, and Urbanovici-Segal macrokinetic models to isothermal crystallization of syndiotactic polypropylene. Thermochimica Acta, 370(1-2), 37-48. http://dx.doi.org/10.1016/S0040-6031(00)00767-X.

8. Zhuomin, D., & Spruiell, J. E. (1997). Interpretation of the nonisothermal crystallization kinetics of polypropylene using a power law nucleation rate function. Journal of Polymer Science. Part B, Polymer Physics, 35(7), 1077-1093. http://

Table 7. Influence of calcium stearate in samples nucleated with NC.

Properties Neat PP PP NC 1 PP NC 1 (without CaSt)

Gloss 45° (%) 54.7 ± 2 63.7 ± 0.9 58.9 ± 0.6Haze (%) 72.5 ± 0.8 45.8 ± 1.1 65.6 ± 0.3Tc (°C) 113 136 133

Figure 12. Haze results of neat and nucleated PP.

Table 6. Influence of calcium stearate in samples nucleated with NB.

Properties Neat PP PP NB 1 PP NB 1 (without CaSt)

Gloss 45° (%) 54.7 ± 2 55.4 ± 1.6 58.4 ± 0.4Haze (%) 72.5 ± 0.8 76.5 ± 2.4 64.8 ± 1.8Tc (°C) 113 132 132

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dx.doi.org/10.1002/(SICI)1099-0488(199705)35:7<1077::AID-POLB7>3.0.CO;2-T.

9. Li, C. C., Zhang, D., & Li, Z. Y. (2002). The effects of alkaline earth dehydroabietate on the crystallization process of polypropylene. Journal of Applied Polymer Science, 85(13), 2644-2651. http://dx.doi.org/10.1002/app.10545.

10. Zhao, X. E., & Dotson, D. L. (2002). US Patent No. 6465551 B1. Alexandria, Virginia: USPTO.

11. Jang, G. S., Cho, W. J., & Ha, C. S. (2001). Crystallization behavior of polypropylene with or without sodium benzoate as a nucleating agent. Journal of Polymer Science. Part B, Polymer Physics, 39(10), 1001-1016. http://dx.doi.org/10.1002/polb.1077.

12. Libster, D., Aserin, A., & Garti, N. (2006). A novel dispersion method comprising a nucleating agent solubilized in a microemulsion, in polymer matrix. I. Dispersion method and polymer characterization. Journal of Colloid and Interface Science, 299(1), 172-181. http://dx.doi.org/10.1016/j.jcis.2006.01.064. PMid:16554065.

13. Botkin, J. H., Dunski, N., & Maeder, D. (2002). Improving molding productivity and enhancing mechanical properties of polypropylene with nucleating agents. Kwinana: Ciba Speciality Chemicals.

14. Santamaria, E., Phan, H. D., & Killough, L. (2008). Clarified polypropylene – old technology vs. new chemistry. In SPE Polyolefins 2008 Proceedings (pp. 1686-1692). Houston: SPE.

15. Wang, K., Zhou, C., Tang, C., Zhang, Q., Du, R., Fu, Q., & Li, L. (2009). Rheologically determined negative influence of increasing nucleating agent content on the crystallization of isotactic polypropylene. Polymer, 50(2), 696-706. http://dx.doi.org/10.1016/j.polymer.2008.11.019.

16. Hobbs, J. K. (2003). In-situ AFM of polymer crystallization. Chinese Journal of Polymer Science, 21(2), 129-133.

17. Tenma, M., & Yamaguchi, M. (2007). Structure and properties of injection-molded polypropylene with sorbitol-based clarifier. Polymer Engineering and Science, 47(9), 1441-1446. http://dx.doi.org/10.1002/pen.20839.

18. Nogales, A., Mitchell, G. R., & Vaughan, A. S. (2003). Anisotropic crystallization in polypropylene induced by deformation of a nucleating agent network. Macromolecules, 36(13), 4898-4906. http://dx.doi.org/10.1021/ma0343028.

19. Kurja, J., & Mehl, N. A. (2001). Plastics additives handbook. Munich: Hanser.

20. Dieckmann, D. (2001). Effect of various acid-neutralizers on the crystallization temperature of nucleated polypropylene. Journal of Vinyl and Additive Technology, 7(1), 51-55. http://dx.doi.org/10.1002/vnl.10264.

Received: May 08, 2015 Revised: Oct. 30, 2015

Accepted: Dec. 15, 2015

http://dx.doi.org/10.1590/0104-1428.2127

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Biodegradation of additive PHBV/PP-co-PE films buried in soil

Barbara Rani-Borges1, Adriano Uemura Faria1, Adriana de Campos2, Suely Patricia Costa Gonçalves1 and Sandra Mara Martins-Franchetti1*

1Polymer Treatment Laboratory – LTP, Biochemistry and Microbiology Department, Universidade Estadual Paulista – UNESP, Rio Claro, SP, Brazil

2National Nanotechnology Laboratory for Agriculture – LNNA, Embrapa Instrumentação – EMBRAPA, São Carlos, SP, Brazil

*sandrafranchetti@hotmail.com

Sbstract

There is considerable concern about the impact plastic materials have on the environment due to their durability and resistance to degradation. The use of pro-oxidant additives in the polymer films could be a viable way to decrease the harmful effects of these discarded materials. In this study, films of PHBV/PP-co-PE (80/20 w/w) and PHBV/PP-co-PE/add (80/19/1 w/w/w) (with pro-oxidant additive) were employed to verify the influence of the additive on the biodegradation of these films in the soil. These films were obtained by melting the pellets in a press at 180 °C which were buried in soil columns for 3 and 6 months. Some samples were also heated before being buried in soil. The biodegradation is higher for the additive blend buried for 3 months than for the pre-heated blend. After 6 months the blend buried and heated/buried was completely degraded in soil. The effect of the additive, on chain oxidation, is more time-dependant than heat-dependant.

Keywords: biodegradation, additive, blend, PHBV/PP-co-PE, soil.

1. Introduction

The increasing use of polymers by our society has resulted in a large quantities of discarded plastic materials in landfills[1,2]. The polymers are employed as raw material to produce different products, such as soft drink bottles, supermarket bags, wall paints, toys, kitchenware, pipes, car dashboards, freezers, tables and many types of daily use items[3]. The inappropriate disposal of these materials is a cause for concern due to the following main characteristics: they are resistant to microbial degradation due to their hydrophobicity and high molecular weight[4-6]. To reduce the impacts of these recalcitrant materials in the environment, researchers are investigating some alternatives, as for example, the production of more degradable polymers, such as the biodegradable and oxo-biodegradable polymers which are less damaging to the environment and production blends composed of synthetic and biodegradable polymers[6,7]. Oxo-biodegradable polymers are synthetic polymers mixed with organic salts (stearate, carbamate) of metals such as cobalt, manganese, iron, etc[8]. These substances cause polymer chain scission in smaller fragments rich in oxygen groups, leading to a decrease in molecular weight and these chains used as carbon source by the microorganisms. According to Scott[9], the oxo-biodegradable polymers undergo a two-step degradation process: abiotic – chemical degradation accelerated by a catalyst, biotic-degradation using microorganisms, which assimilate the oxidation products.

The principle of oxo-biodegradable polymers is based on the addition of functional chemical groups to the polymer chain, which enables the material to disintegrate, facilitating the subsequent microbial biodegradation process[10,11].

Another alternative to minimize the damage caused by plastic disposal is implementing pre-treatment methods as UV radiation and heat to facilitate the microbial adhesion in the biodegradation process[12,13].

The most widely used synthetic polymers are polyolefins: PP, PE, which when discarded in the environment have long term durability. These polymers can be mixed (blended) with a biodegradable polymer as poly(hydroxybutirate) (PHB) or poly(hydroxybutirate)-co-valerate (PHBV) and with pro-oxidant additives to make the material more degradable.

PP and PE are mechanically resistant, non biodegradable thermoplastic resins, widely used in different applications and produce large volumes of waste because of their resistance to microbial attack. These polymers are susceptible to UV light, oxidation and temperature[14]. Notwithstanding these limitations, they are economically viable cause produce a copolymer of PP and PE (PP-co-PE)[8,15], a new material with more flexibility, processability and tenacity[16].

The interest on polymer blends has grown recently due to the fact these materials match different properties originating new materials with good mechanical, chemical and thermal properties. Moreover, the cost to produce blends is less than the cost to produce new polymers[17-20]. In this work we used the blend PHBV/PP-co-PE (80:20 w/w), with a biodegradable and the synthetic copolymer, respectively.

Carashi et al.[19] describe PHB as a natural polymer, produced by some types of bacteria in adequate culture media, and with good resistance to water steam and storage stability for shelf life of at least four years. It is a semi-crystalline

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polymer, very brittle, and this characteristic is improved when the PHBV copolymer is used[20].

According to the American Society of Tests and Materials (ASTM)[21] and the International Organization of Normalization (ISO)[22] degradation is an irreversible process that results in the structural modification of the polymer, which causes the loss of basic properties. This process is directly influenced by the environmental conditions, such as pH, temperature, humidity, sun light and microbial composition[23], besides the material structure, its surface area and morphology[24]. Biodegradation involves the microorganisms using the polymer as the sole carbon source, which then generate biofilm on the polymer surface, containing water and extracellular polymer substances that cause damage and scission of the macromolecular chain in simple molecules[25]. If the biodegradation is complete – mineralization – these simple molecules degrade in CO2, water and cellular biomass, in aerobic conditions or CO2, CH4 and biomass in anaerobic environment[26]. The breaking of polymeric bonds is related to the enzyme activity of the microorganisms in two ways: outside or inside of cells, i.e., exo or endo-biodegradation, respectively, or through a combination of these two ways[26]. The two processes are mediated by enzymes: the hydrolases and oxidases. The first process can break down the polymers by adding water molecules in the polymer chains, and the second one inserts oxygen (peroxide groups) in the chains which are also broken down[20,27].

Biodegradation depends on the polymer composition, molecular weight, crystallinity, presence of functional groups, type of microorganisms[27-29], therefore there are differences in the degradation stages that different polymers undergo: PHBV, PP/PE and polymers containing pro-oxidant additives:

➢ PHBV: this polymer contains hydrophilic groups able to interact with water molecules. The hydrolases convert the polymer into carboxyl acid, whereas the extracellular degradation produces 3-hydroxybutirate and 3-hydroxyvalerate molecules, which are metabolized inside cells, producing carbon dioxide and water[28].

➢ PP/PE: These polymers are considered inert due to their hydrophobic characteristics such as high molecular weight and lack of functional groups to facilitate the microbial attack[25].

➢ Additive polymer: the oxi-biodegradable plastics undergo two degradation stages, abiotic and biotic. In the first stage there is a reduction in molecular weight by the oxidation chains, creating carboxylic acids, alcohols and ketones. This stage allows the second step: the chains become more hydrophilic favoring extracellular enzyme activity and subsequent breakdown of polymer chains[30,31].

Thus, the objective of this work is to investigate the biodegradation of poly(hydroxybutyrate-co-hidroxyvalerate) (PHBV) blended with polypropylene-co-polyethylene (PP-co-PE), with and without a pro-oxidant additive, through biotreatment in soil columns, using weight loss, scanning electronic microscopy (SEM), Fourier transform infrared (FTIR) and X-ray diffraction (XRD). There are different factors that affect biodegradability and to the best of our knowledge, there are a few papers that address biodegradation of PHBV blended with PP-co-PE, with and

without a pro-oxidant additive. The effect of the additive depends on the biodegradation period and not on the heat pre-treatment.

2. Materials and Methods

2.1 Materials

PHBV with 6.2% of HV and Mϖ650.000 – was supplied by PHB Industrial SA, g/mol. PP-co-PE Mw 220.000 g/mol was supplied by Braskem. The pro-oxidant additive was supplied by RES Brazil LTDA, stated as manganese dithiocarbamate, impregnated in polyethylene matrix.

2.2 Methods2.2.1 Films preparation

The blends PHBV/PP-co-PE (80/20w/w) and PHBV/PP-co-PE/add (80/19/1 w/w/w) with pro-oxidant additive were prepared using an internal mixer (Haake Torque Rheometer), 50 rpm, 180 °C for 5 minutes. The blend composition chosen was to investigate how the synthetic copolymer (lower quantity) influences the material degradation. The materials were pressed at 170 °C under 71.3 kgfcm–2 for 3 min to produce films of 5 cm diameter and 80-100 μm thickness. The films (in duplicate) were sterilized before being placed in soil columns, with 2% solution of sodium hypochlorite.

2.2.2 Treatments

2.2.2.1 Heat

The films were heated in a vacuum oven (Marconi – MA 030), at 100 °C, for 40 h. The heat treatment was applied for 5 h/day. The films were slowly cooled down inside the oven. Next, they were buried in soil columns.

2.2.2.2 Buried-soil columns

The soil columns were built using 1.5 L of PET bottles with small holes to aerate the soil during the experiment. The the bottle was partially filled with soil, 12 cm at the bottom of the bottle covered with the film, then more soil, 8 cm, was placed over the film (Figure 1). The soil used was collected from the campus of Rio Claro, SP, Brazil, rich in humus, without litter and leaves, sieved with a 2 mm mesh, and initially adjusted to have 60% of humidity[32]. The humidity was maintained by a water recipient placed under the bottle, allowing the water to rise through capillarity, described by Campos et al.[33] and shown in Figure 1. The columns were kept at room temperature for 3 months and 6 months (March and November /2012).

Figure 1. Soil column.

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2.2.3 Characterization of the polymers

2.2.3.1 Weight loss

The weight loss was evaluated by the difference in mass before and after the biotreatment, weighed on an analytical balance, CHYO, model JK200.

2.2.3.2 Fourier Transform Infrared (FTIR)

The original (before biological treatment) and biotreated films were analyzed by a Shimadzu IR Prestige 21 FTIR spectrometer, 16 scans with 4 cm-1 resolution, and range of 4000 to 400 cm–1. To compare the spectra before and after the soil treatment, each blend spectrum was normalized by the internal standard at 1460 cm–1, assigned to vibration of the CH2groups, because this band does not change during the biodegradation process.

2.2.3.3 X-Ray Diffraction (XRD)

The samples were analyzed by XRD in a Shimadzu LABX-XRD 6000 diffractometer, operating at 30 kV, 30 mA and CuKα (λ 1.5406 Å). The assays were performed at 25 °C, with 2ϴ angles at 10 and 35° (2º/min). Some samples were completely deteriorated after the biodegradation and were not measured.

The crystallinity degrees (Xc) were calculated by the ratio of the crystalline peak and the total area (crystalline + amorphous peaks), using the peak deconvolution method with Origin 7.5 software, using Gauss function to define the shape of the peaks, after the baseline correction.

2.2.3.4 Scanning Electron Microscopy (SEM)

SEM was used to investigate the surface morphology of the polymer in order to elucidate the material properties of polymers and behavior of some processes such as: thermal, photo-degradation and biodegradation[34].

The polymer films before and after the heat and microbial treatments were analyzed by a scanning electron microscope - Zeiss DSM 940-A, using acceleration voltage of 5 kV (ESALQ/USP – Piracicaba, SP).

3. Results and Discussions

After 3 and 6 months the PHBV/PP-co-PE films with and without additive were analyzed to verify morphological and structural changes, using the methods cited above.

There was a significant weight loss in the blends, which was higher in the PHBV/PP-co-PE/Add blend (Table 1). The additive (pro-oxidant), favored the chains scission and biodegradation, but the presence of PHBV was also more significant to biodegradation.

After 6 months, the blends with additive were not found in the soil columns, just some small pieces of the blend without additive, thus the weight loss was considered to be 100%.

FTIR measurements compare the spectra of the films before and after the heat and soil treatment (Figure 2). These spectra were normalized by an internal standard band (here called A0) at 1460 cm–1 (assigned to the PHBV CH2 deformation, as this band does not change after treatment) and deconvolution applied by the Lorentzian function to adjust the curves and isolate each band with the corresponding area, from the overlapping bands, increasing the spectrum resolution[35].

Table 2 shows a comparison of carbonyl indices in the amorphous and crystalline phases from the original and treated spectra, which were calculated by the areas ratio: AC=O/A0 bands. In the PHBV/ PP-co-PE blend there was a significant decrease in C=O indices, in both phases: amorphous and crystalline (51% and 57%, respectively), emphasizing that the blend with additive underwent more degradation (83% and 92%, respectively), when compared with the neat samples. In the heated and buried samples there was a synergistic effect in the amorphous phase and the decrease in the C=O was 81%. Considering the two subsequent treatments, heat/buried, the additive samples showed the heat effect on the chain structure, before the soil treatment. These results showed that the action of the oxidant additive in the soil biodegradation process was very efficient, more efficient than the heat/buried degradation, leading to a greater scission of polymer chains, especially in the crystalline phase. This fact suggests that the less ordered valerate chains of the PHBV located at the edge of the PHB crystal nucleus favors the microbial action, as in the model proposed by Yoshie[36]. Sadi et al.[37] observed that PHB samples exposed to UV radiation caused several changes in the material, such as scission and crosslinking reactions. An initial delay in the biodegradation was also observed, which was related to a thin superficial layer with higher crystallinity of the samples exposed to UV radiation. However, this did not completely inhibit PHB decomposition and as this layer was consumed there was an improvement in the biodegradation due to the fact that the degraded molecules located underneath this layer did not reorganize themselves into crystals.

The PHBV/ PP-co-PE blends buried for 6 months were measured using the FTIR technique (using very small pieces) and the graphics (in Figure 2) showed that the remaining parts

Table 1. The weight loss of the samples.Samples Buried Heat/Buried

PHBV 90.1 100PP-co-PE 0.12 0.23PHBV/ PP-co-PE 77.0 93.0PHBV/PP-co-PE/Add 88.0 94.0

Table 2. Comparison carbonyl indices in the amorphous and crystalline phases from the original and treated spectra of blends with and without additive.

Samples Neat Buried Heat/Buried

PHBV/PP-co-PEAmorphous phase (A 1727/1460) 0.75 0.37 0.14Crystalline phase (A1716/1460) 0.83 0.36 0.35

PHBV/PP-co-PE/AddAmorphous phase (A1751/1460) 0.66 0.11 0.10Crystalline phase (A1715/1460) 0.86 0.07 0.35

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in the samples were principally constituted by the synthetic copolymer, thus the PHBV was almost completely consumed (Figure 2a, b). PHBV/ PP-co-PE and PHBV/ PP-co-PE/add (heat/buried) were not found in the soil. In this period the additive blend buried in the soil also disappeared. These samples were considered totally biodegraded.

Table 3 shows the crystallinity degrees (Xc) of the polymer films before and after the heat treatment and Figure 3a, b displays the diffractograms. The heat treatment caused a crystallinity increase of the blend (4%) and additive blend (12%), causing the annealing. Comparing these two blends, the crystallinity increased more in the additive

blend than in the blend without additive, due to the fact that more chains break and reorganize. PP-co-PE blend films after heat/microbial treatment were not measured by XRD because they were too deteriorated.

Figure 4 shows the micrographs of PHBV/ PP-co-PE and PHBV/ PP-co-PE/add. These figures clearly show biodegradation on the surface of the samples, without and with additive, such as microbial adhesion and surface delamination of PHBV. The morphological changes in the blends with additive seem to be greater than in those without additive, according to the visual aspect of these damaged samples.

Figure 2. FTIR measurements comparison of the films before and after the heat and soil treatment.

Table 3. Crystallinity degrees (Xc) of the polymer films before and after the heat treatment.Samples hkl 2θ (°) D (nm) % Xc

PHBV/PP-co-PE-neat(020) PHBV 12.9 2.38

81(110)PHBV (040)PP 16.4 1.90

PHBV/PP-co-PE-heat(020) PHBV 13.7 2.52

85(110)PHBV (040)PP 17.0 1.91

PHBV/PP-co-PE/Add – neat(020) PHBV 12.9 2.37

69(110)PHBV (040)PP 16.4 1.90

PHBV/PP-co-PE/Add –heat(020) PHBV 12.9 2.86

81(110)PHBV (040)PP 16.4 2.19

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Figure 3. (a) Diffractograms of the PHBV/PP-co-PE films before and after the heat treatment; (b) Diffractograms of the PHBV/PP-co-PE/Add films before and after the heat treatment.

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Polímeros, 26(2), 161-167, 2016166

Figure 4. Micrograph of PHBV/ PP-co-PE (a, b, c) and PHBV/ PP-co-PE/add (d, e, f).

4. Conclusions

The PHBV/PP-co-PE blend heat/buried in soil for 3 months underwent significant degradation. In the buried additive blend the biodegradation was higher, i.e., the additive assisted the oxidation chains.

The PHBV/PP-co-PE blend heat/buried was completely degraded after 6 months. Moreover, the buried additive blend was mineralized in the soil. The action mechanism of the additive was not heat-dependant, showing that time was more influential.

5. Acknowledgements

The authors are grateful to Fapesp, CNPq and Capes for the grants, scholarship, and financial support.

The authors are also grateful to Dr. Marco Aurelio De Paoli from the Chemistry Institute, Unicamp, Campinas, SP, Brazil for his support in using the Haake Torque Rheometer and Dr. José A. M. Agnelli from Department of Materials Engineering, São Carlos Federal University, São Carlos, SP, Brazil for supplying the PHBV.

6. References

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PE-PHBV blends. International Journal of Material Science, 3(2), 54-60.

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10. Chiellini, E., Corti, A., & Swift, G. (2003). Biodegradation of thermally-oxidized, fragmented low-density polyethylenes. Polymer Degradation & Stability, 81(2), 341-351. http://dx.doi.org/10.1016/S0141-3910(03)00105-8.

11. Fechine, G. J. M., Santos, J. A. B., & Rabello, M. S. (2006). Avaliação da fotodegradação de poliolefinas através de exposição natural e artificial. Quimica Nova, 29(4), 674-680. http://dx.doi.org/10.1590/S0100-40422006000400009.

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12. Guedri, L., Amor, S. B., Gardettec, J. L., Jacqueta, M., & Rivaton, A. (2005). Lifetime improvement of poly(ethylene naphthalate) by ZnO adhesive coatings. Polymer Degradation & Stability, 88(2), 199-205. http://dx.doi.org/10.1016/j.polymdegradstab.2004.05.015.

13. Wilhelm, H. M., & Felisberti, M. I. J. (2002). Reactive compatibilization of maleated polypropylene and maleatedpoly(styrene-b-butadiene-b-styrene) blends. Journal of Applied Polymer Science, 86(2), 366-371. http://dx.doi.org/10.1002/app.10977.

14. Coutinho, F. M. B., Mello, I. L., & Santa Maria, L. C. (2003). Polietileno: Principais tipos, propriedades e aplicações. Polímeros: Ciência e Tecnologia, 13(1), 1-13. http://dx.doi.org/10.1590/S0104-14282003000100005.

15. Fernandes, L. L., Freitas, C. A., Demarquette, N. R., & Fechine, G. J. M. (2012). Estudo do efeito do tipo de polipropileno na fotodegradação da blenda polipropileno/poliestireno de alto impacto. Polímeros: Ciência e Tecnologia, 22(1), 61-68. http://dx.doi.org/10.1590/S0104-14282012005000013.

16. Silva, D. F., Araújo, E. M., & Melo, T. J. A. (2012). Desenvolvimento de blendas de poliamida 6/composto de borracha reciclada (SBR-R). Revista Eletrônica de Materiais e Processos, 7(2), 104-110.

17. Chen, J. H., Zhong, J. C., Cai, Y. H., Su, W. B., & Yang, Y. B. (2007). Morphology and thermal properties in the binary blends of polypropylene-co-ethylene copolymer and isotatic polypropylene with polyethylene. Polymer, 48(10), 2946-2957. http://dx.doi.org/10.1016/j.polymer.2007.03.037.

18. Pospíšil, J., Horák, Z., Kruliš, Z., Nešpůrek, S., & Kuroda, S. (1999). Degradation and aging of polymers blends. Polymer Degradation & Stability, 65, 405-414. http://dx.doi.org/10.1016/S0141-3910(99)00029-4.

19. Carashi, J. C., Ramos, U. M., & Leão, A. L. (2002). Compósitos biodegradáveis de polihidroxibutirato (PHB) reforçado com farinha de madeira: propriedades e degradação. Acta Scientiarum, 24(6), 1609-1614. http://dx.doi.org/10.4025/actascitechnol.v24i0.2475.

20. Franchetti, S. M. M., & Marconato, J. C. (2006). Polímeros biodegradáveis: uma solução parcial para diminuir a quantidade dos resíduos plásticos. Quimica Nova, 29(4), 811-816. http://dx.doi.org/10.1590/S0100-40422006000400031.

21. American Society for Testing and Materials – ASTM. (1991). ASTM D883: standard terminology relating to plastics. West Conshohocken: ASTM.

22. International Organization for Standardization – ISO. (1997) ISO 846: plastics: determination of behavior under the action of fungi and bacteria: evaluation of the action of microorganisms. Geneva: ISO.

23. Imam, S. H., Gordon, S. H., Shogren, R. L., Tosteson, T. R., Govind, N. S., & Greene, R. V. (1999). Degradation of starch PHBV bioplastic. Applied and Environmental Microbiology, 65, 431-437. PMid:9925564.

24. Chandra, R., & Rustgi, R. (1998). Biodegradable polymers. Progress in Polymer Science, 23(7), 1273-1335. http://dx.doi.org/10.1016/S0079-6700(97)00039-7.

25. Arkatkar, A., Arutchelvi, J., Bhaduri, S., Uppara, P. V., & Doble, M. (2009). Degradation of unpretreated and thermally

pretreated polypropylene by soil consortia. International Biodeterioration & Biodegradation, 63(1), 106-111. http://dx.doi.org/10.1016/j.ibiod.2008.06.005.

26. Krzan, A., Hemjinda, S., Miertus, S., Corti, A., & Chiellini, E. (2006). Standardization and certification in the area of environmentally degradable plastics. Polymer Degradation & Stability, 91(12), 2819-2833. http://dx.doi.org/10.1016/j.polymdegradstab.2006.04.034.

27. Lucas, N., Christophe, B., Belloy, C., Queneudec, M., Silvestre, F., & Nava-Saucedo, J. E. (2008). Polymer biodegradation: mechanisms and estimation techniques. Chemosphere, 73(4), 429-442. http://dx.doi.org/10.1016/j.chemosphere.2008.06.064. PMid:18723204.

28. Shah, A. A., Hasan, F., Hameed, A., & Ahmed, S. (2008). Biological degradation of plastics: a comprehensive review. Biotechnology Advances, 26(3), 246-265. http://dx.doi.org/10.1016/j.biotechadv.2007.12.005. PMid:18337047.

29. Priyanka, N., & Aschana, T. (2011). Biodegradability of Polythene and plastics by the help of microorganism: A way for brighter future. Environmental and Analytical Toxicology, 1, 1-4. http://dx.doi.org/10.4172/2161-0525.1000111.

30. Brito, G. F., Agrawal, P., Araújo, M., & Melo, T. J. A. (2011). Biopolímeros, polímeros biodegradáveis e polímeros verdes. Revista Eletrônica de Materiais e Processos, 6(2), 127-139.

31. De Paoli, M. A. (2008). Degradação e estabilização de polímeros. São Paulo: Editora Artliber.

32. Campos, A., Marconato, J. C., & Martins-Franchetti, S. M. (2011). Biodegradation of blend films PVA/PVC, PVA/PCL in soil and soil with landfill leachate. Brazilian Archives of Biology and Technology, 54(6), 1367-1377. http://dx.doi.org/10.1590/S1516-89132011000600024.

33. Campos, A., Marconato, J. C., & Martins-Franchetti, S. M. (2010). Biodegradação de Filmes de PP/PCL em Solo e Solo com Chorume. Polímeros: Ciência e Tecnologia, 20(4), 295-300. http://dx.doi.org/10.1590/S0104-14282010005000039.

34. Kaczmarek, H. (1996). Changes of polymer morphology caused by U.V. irradiation: surface damage polymer. Polymer Degradation & Stability, 37, 189-194. http://dx.doi.org/10.1016/0032-3861(96)81086-X.

35. Araújo, S. C., & Kawano, Y. (2001). Espectro Vibracional no infravermelho próximo dos polímeros poliestireno, poli(Metacrilato de Metila) e policarbonato. Polímeros: Ciência e Tecnologia, 11(4), 213-221. http://dx.doi.org/10.1590/S0104-14282001000400011.

36. Yoshie, N., Saito, M., & Inoue, Y. (2001). Structural Transition of lamella crystals in a isomorphous copolymer, poly(3-hydroxybutyrate-co-3-hydroxyvalerate). Macromolecules, 34(26), 8953-8960. http://dx.doi.org/10.1021/ma0113071.

37. Sadi, R. K., Fechine, G. J. M., & Demarquette, N. R. (2010). Photodegradation of poly(3-hydroxybutyrate). Polymer Degradation & Stability, 95(12), 2318-2327. http://dx.doi.org/10.1016/j.polymdegradstab.2010.09.003.

Received: Mar. 11, 2015 Revised: Aug. 23, 2015

Accepted: Dec. 15, 2015

http://dx.doi.org/10.1590/0104-1428.2013

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The effect of andiroba oil and chitosan concentration on the physical properties of chitosan emulsion film

Vanessa Tiemi Kimura1, Cintia Satiyo Miyasato1, Bianca Pereira Genesi1, Patrícia Santos Lopes1, Cristiana Maria Pedroso Yoshida1 and Classius Ferreira da Silva1*

1Laboratório de Biotecnologia e Produtos Naturais – BIONAT, Instituto de Ciências Químicas, Ambientais e Farmacêuticas, Universidade Federal de São Paulo – UNIFESP, Diadema, SP, Brazil

*cfsilva@unifesp.br

Sbstract

Chitosan film is used as a dressing to heal burns. The physical and biological properties of the film can be modified by the addition of phytotherapic compounds. This work used the casting -solvent evaporation technique to prepare chitosan film containing andiroba oil (Carapa guianensis) which has anti-inflammatory, antibiotic, and healing properties. The objective of this study was to determine the effect of the concentrations of chitosan and andiroba oil on the physical properties of chitosan films. The emulsion films were evaluated concerning the mechanical properties and fluid handling capacity. Additionally, scanning electron microscopy and thermal analysis were performed. The results showed that the barrier and mechanical properties were affected by the addition of andiroba oil, and these may be modulated as a function of the concentration of oil added to the film. The thermal analysis showed no evidence of chemical interactions between the oil and chitosan.

Keywords: biopolymers, dressings, Carapa guianensis.

1. Introduction

Andiroba (Carapa guianensis) is a tree of the Meliaceae family found in Central and South America, especially in the Amazon basin region. Many studies have reported the pharmacological properties of products obtained from the andiroba flower[1,2], from the ethanolic extract of the andiroba leaf[3,4] and especially products derived from andiroba seed oil[5-7]. Andiroba oil is widely used in popular medicine in the Amazon basin region.

According to Cabral et al.[8], andiroba oil is composed mostly of triacylglycerols, with high levels of unsaturated and saturated fatty acids such as oleic (51.81%), palmitic (25.76%), stearic (9.08%), and linoleic (8.3%).The medicinal properties of andiroba oil have been attributed to the presence of limonoids, which are tetranortriterpenoids[6]. Andiroba oil also contains triterpenes, tetraterpenes, alkaloids, and glycerides[9].

Andiroba seed oil is currently considered to be acaricides[10], larvicidal against Aedes aegypti[11], as well as being an antiplasmodial[5], anti-inflammatory[12], anti-allergic[7,13,14], and also suitable for wound healing[3,4].

The ethanolic extract of Carapa guianensis leaves was evaluated for antibacterial and wound healing activity, using excision, incision, and dead space wound models in rats. The results showed an increased rate of wound contraction and hydroxyproline content (a biochemical marker for tissue collagen), which indicates the potential application of Carapa guianensis in wound healing[4].

Many wound dressings have been developed for the treatment of severe burn wounds or ulcers. Damaged tissue requires biocompatible materials like chitosan that has a high film-forming capacity. The most cited advantages of chitosan are its physico-chemical and biological

properties. Chitosan promotes activation and proliferation of inflammatory cells in granular tissues[15], stimulates cell proliferation and histoarchitectural reorganization of the tissue[16], and affects the functioning of macrophages, thus accelerating the healing process[17]. The use of chitosan resulted in a substantial decrease in healing time and minimal scarring in several animals[18]. These and other properties can be potentiated with the incorporation of andiroba oil. As previously described, andiroba oil has many interesting properties for use in dressings.

The preparation of emulsified films presents some challenges, for example, maintaining the stability of the emulsion during the process of drying the film. The stability of chitosan films can be achieved by adding surfactants but also by emulsification with rotor–stator homogenizer (> 20,000 rpm)[19-21], or even less vigorously (13,500 rpm)[22]. The small concentration of chitosan in oil emulsions also helps to maintain the stability of the emulsion. Usually, this concentration ranges from 0.1 to 1%[19-21], but it can reach values up to 3%[22]. Therefore, in this work, we decided to work with concentrations between 0.1 and 1% and agitation (24,000 rpm).

This work aims to study the effect that the concentrations of andiroba oil and chitosan have on the physical and chemical properties of chitosan film used for wound dressings.

2. Materials and Methods

2.1 Materials

Commercial chitosan (deacetylation of approximately 82% and molar mass of approximately 1.47 × 105 g/mol) was supplied by Polymar (Fortaleza, Brazil) without prior

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Polímeros, 26(2), 168-175, 2016 169

purification. Acetic acid (Synth, Brazil) was used as an acidic medium. Andiroba oil was provided by MPR Indústria e Comércio de Óleos Vegetais Ltda (Brazil).

2.2 Chitosan suspension

Chitosan (1.0% or 2.0%, w/w) was dissolved in an aqueous acetic acid solution. The stoichiometric amount of acetic acid was calculated to achieve the protonation of all the NH2 sites and taking into account the sample weight and the degree of acetylation. Using this base amount plus an extra 50% gave the total stoichiometric amount to be used. The suspension was homogenized for 2 h prior to the preparation of the chitosan film, in order to complete chitosan solubilization.

2.3 Preparation of chitosan film

The films were prepared by the casting technique. The chitosan suspension and the andiroba oil was emulsified beforehand (24000 rpm for 10 min). We tested three concentrations of andiroba oil (g of andiroba oil/100 g of solution, i.e. % w/w): 0.1%, 0.5%, and 1.0% w/w. The chitosan emulsion was poured into polyethylene Petri dishes. The films were dried in a forced air oven at 40 °C for 24 h. The mass of the suspension applied to the Petri dishes was kept constant (0.21 g/cm2).The films then underwent various analyses.

2.4 Scanning Electron Microscopy (SEM)

SEM analysis was performed on fractured cross-sections and the surfaces of gold-sputtered films using an LEO 440i scanning electron microscope (LEO Electron Microscopy Ltda.) with 10 kV and 100 pcA.

2.5 Fluid Handling Capacity (FHC)

The fluid handling capacity (FHC) of the film is defined as the sum of the Absorbency (ABS) and Moisture Vapor Transmission Rate (MVTR). The FHC was examined according to the BS EN 13726-1 method[23] for hydrocolloids and dressings. In this test, samples of each film (or dressing) were applied to the modified Paddington cups (Figure 1), to which were added 20 mL of simulated exudate fluid (SEF).

The cups were weighed using a calibrated analytical balance, inverted so that the dressing came into contact with

the SEF – see Figure 1c – and the solution was placed in a temperature and humidity controlled incubator to maintain an environment of 37 °C ± 2 and a relative humidity below 20% for a period of 24 h. At the end of the test the cups were removed from the incubator and were allowed to equilibrate at room temperature for a period of 30 min prior to reweighing on the analytical balance. The FHC, ABS, and MVTR were calculated by the following equations:

x yMVTR

time surface−

=×

(1)

b aABStime surface

−=

× (2)

FHC MVTR ABS= + (3)

where x is the complete system weight (film + SEF solution + cup) at the beginning of the test; y is the complete system weight (film + SEF solution + cup) after 24 h; b is the film weight at the beginning of the test; and a is the film weight after 24 h. Five repetitions were done per experiment.

2.6 Mechanical properties

Tensile testing was done in accordance with the ASTM D882 method[24]. Films were cut into 10.00 cm × 2.54 cm strips. The tensile strength, elongation at breaking point, and Young’s modulus were measured using TexturePro CT V1.2 (Brookfield, CT3 50K Texturometer). The crosshead speed was set at 1 mm.s–1. Samples were pre-conditioned in a desiccator at 75% relative humidity for 48 hours. There were at least 10 repetitions per experiment.

2.7 Thermal analysis

Differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA) studies were performed on chitosan film, pure chitosan, and andiroba oil. TGA was done with a TGA-60 (Shimadzu) analyzer. All analyses were performed with 5-10 mg samples in platinum pans in a dynamic nitrogen atmosphere (100 mL.min–1), between 30 °C and 700 °C. The experiments were done at a scanning rate of 10 °C.min–1. DSC analysis was performed with a DSC-60 (Shimadzu) analyzer. Samples (approx. 5-10 mg) were scanned in a sealed aluminum pan and heated to a

Figure 1. Modified Paddington cups used for the determination of FHC.

Kimura, V. T., Miyasato, C. S., Genesi, B. P., Lopes, P. S., Yoshida, C. M. P., & Silva, C. F.

Polímeros, 26(2), 168-175, 2016170

temperature of 450 °C at a rate of 10 °C.min–1, in a nitrogen atmosphere, and with a flow rate of 50 mL.min–1.

2.8 Statistical analysis

All the characterizations were done in replicate. The Tukey’s test was done for comparison of means, using BioEstat 5.3[25].

3. Results and Discussion

3.1 Scanning electron microscopy

No macroscopic phase separation was observed after drying for any of the samples prepared from chitosan/andiroba oil, which indicates that oil droplets were stabilized in the chitosan suspension. SEM was used to evaluate the morphology and distribution of oil droplets in the films. This analysis allows us a better understanding of mechanical and barrier properties[26]. Figure 2 shows SEM micrographs of the chitosan

films containing andiroba oil for 2% chitosan. Micrographs of films for 1% (w/w) chitosan were quite similar, and they are not presented. The emulsified films showed structural discontinuities associated with the formation of two phases (lipid and polymer) in the matrix. The oil-free films had a smooth and homogeneous microstructure with no irregularities like air bubbles or oil droplets detected (micrograph not shown), as our group had previously published[27]. The number of oil droplets increases as the concentration of oil increases. The cross-sections of the emulsified films show that droplets have an ellipsoidal shape, which was also verified by other authors[26,28,29]. This ellipsoidal shape can be attributed to the weight of the chitosan over the droplets during the drying. It is important to mention that no surfactant was used in the preparation of our films.

Kokoszka et al.[30] studied whey protein/rapeseed oil emulsion film.Unlike what was observed in this study, they verified that the oil is not well distributed throughout

Figure 2. SEM of the surface (left) of chitosan film viewed at a magnification of 500×, and cross-sections (right) viewed at a magnification of 3000× (Chitosan 2.0% w/w).

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Polímeros, 26(2), 168-175, 2016 171

the film on both sides. According to Kokoszka et al.[30], the oil droplets are more concentrated on the side exposed to the air since the film retraction during drying produces changes in its structure, which becomes denser, and the oil droplets migrate towards the side exposed to the air, thus favoring coalescence.

3.2 Fluid handling capacity

The moisture content of the wounds must be carefully controlled to achieve optimal rates of wound healing. The healing process can be influenced by changes in the moisture content of a wound and the skin around it. A wound that is too dry may delay or impair the healing while excess fluid can cause maceration or infection. Thus, the optimal healing environment is achieved by applying an appropriate dressing that should be removed in time to avoid maceration or adherence[31]. Donor sites, unspecified granulating wounds, and third-degree burns generate between 3.4 and 5.1g of exudate per 10 cm2 over a 24 hour period[32].

The MVTR, ABS, and FHC are presented in Table 1. Thomas and Young[31] evaluated these properties for two commercial dressings: ActivHeal (Advanced Medical Solutions) and Allevyn Adhesive (Smith & Nephew) – both are film-backed foam dressings. They verified the absorbencies to be 3.44 and 4.32 g/10 cm2/24 h, respectively, for ActivHeal and Allevyn Adhesive. Our results for absorbency are very close to these values, especially when the concentration of chitosan is 2.0% w/w. Furthermore, the MVTR were 1.67 and 12.35 g/10 cm2/24 h, respectively, for ActivHeal

and Allevyn Adhesive[31]. Regarding the MVTR, our results were close to those for ActivHeal. Also, it is important to mention that both these commercial dressings are foam dressings, so these properties are usually greater than those for film dressings.

The WVTR decreases with the oil concentration in both chitosan concentrations; this could be attributed to the hydrophobicity of the andiroba oil. Concerning the Absorbency, the incorporation of andiroba oil decreased this parameter especially for the 2% chitosan films.

3.3 Mechanical properties

To adequately protect a wound, the film must maintain its integrity against external stress during the manipulation and application, even when it is on the wound. Tensile strength indicates the maximum tensile stress that the film can sustain, elongation at breaking point is the maximum change in length of a test specimen before breaking, and the Young’s modulus is a measure of the stiffness of the film[21]. The mechanical properties of chitosan film are shown in Table 2. Both the concentrations significantly affected the mechanical properties of the emulsified films.

Young’s modulus and the tensile strength of the chitosan film increased when andiroba oil was incorporated into the chitosan matrix. These two properties are usually higher for samples with 2.0% of chitosan. The addition of low concentrations (0.1%) of andiroba oil causes an increase in these properties if compared to oil-free film. Whenever the concentration of oil increases, these two properties decrease

Table 1. Fluid handling properties of the film dressing for different concentrations of andiroba oil.

Concentration (% w/w)MVTR (g/10cm2/24h) Absorbency

(g/10cm2/24h) FHC (g/10cm2/24h)Chitosan Andiroba oil

1.0

0.0 1.73 ± 0.07ª 1.67 ± 0.12ª 3.39 ± 0.140.1 2.83 ± 0.28b 0.62 ± 0.06b 3.45 ± 0.290.5 1.53 ± 0.12c 1.59 ± 0.16ª 3.12 ± 0.201.0 1.31 ± 0.13d 1.49 ± 0.08ª 2.81 ± 0.15

2.0

0.0 3.24 ± 0.22e 3.43 ± 0.21c 6.67 ± 0.300.1 1.88 ± 0.12ª 2.32 ± 0.18d 4.20 ± 0.210.5 1.28 ± 0.13d 2.74 ± 0.26e 4.02 ± 0.291.0 1.36 ± 0.07c,d 2.63 ± 0.21d,e 3.98 ± 0.22

Different superscripts within the same column indicate significant differences among formulations (p<0.05).

Table 2. Effect of the concentrations of chitosan and andiroba oil on Young’s modulus, tensile strength, and elongation at breaking point, for the emulsified films.

Concentration (% w/w)Young’s modulus (MPa) Tensile strength (N/m2) Elongation at breaking

point (%)Chitosan Andiroba oil

1.0

0.0 36.52 ± 1.49a 169.99 ± 14.06a 13.59 ± 1.20a

0.1 67.67 ± 5.32b 263.47 ± 22.62b 13.48 ±1.29a

0.5 55.39 ± 4.96c 204.16 ± 18.71c 13.07 ± 0.40a

1.0 34.63 ± 3.99a 172.60 ± 17.16a 13.69 ± 1.03a

2.0

0.0 62.02 ± 6.75b,c 171.21 ± 14.01a 22.56 ± 1.86b

0.1 98.50 ± 9.58d 272.55 ± 17.00b 9.26 ± 0.89c

0.5 81.98 ± 3.55e 250.20 ± 22.77b 9.23 ± 0.92c

1.0 57.32 ± 4.37c 182.67 ± 17.11a,c 12.43 ± 1.14a

Different superscripts within the same column indicate significant differences among formulations (p<0.05).

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Polímeros, 26(2), 168-175, 2016172

and are closer to those of the control film (oil-free chitosan film). A decrease in these two properties with the increase of oil was also observed in emulsified films with chitosan/basil essential oil[19], chitosan/thyme essential oil[20], chitosan/tea tree oil[22], cassava starch-chitosan/oregano essential oil[33], and whey protein/olive oil[34].

The mechanical properties obtained by the addition of oil may be related to the structural arrangement of the lipid phase in the chitosan matrix. A number of discontinuities increases as the concentration of andiroba oil increases, which could explain the decrease in Young’s modulus and tensile strength.

No significant effect on elongation at breaking point was observed for 1% of chitosan and when the andiroba oil concentration increased. This effect was also reported by other authors when adding oil to a chitosan matrix[29,35,36] and it could also be attributed to the structural discontinuities provoked by the incorporation of the oil. Moreover, the incorporation of oil promoted a substantial reduction in the elongation of the films containing 2% chitosan.

The effects of chitosan concentration on the mechanical and physical properties of the films seem to be greater than the effect of the oil. Most likely, the oil concentrations are not sufficiently high so that they could be more important than chitosan concentration. But it would be difficult to produce films with higher oil concentration because we would have exudation of the oil from the film, even if we had used one tensoative.

3.4 Thermogravimetric analysis

TGA was performed to evaluate the thermal stability of the chitosan powder, andiroba oil, and chitosan-andiroba oil films. Thermal degradation is displayed in Table 3 and Figures 3, 4.

Table 3 and Figure 3 show that chitosan powder mainly loses mass due to decomposition between 200 °C and 400 °C, and especially in the range of 200 °C to 300 °C. On the other hand, the weight loss of the pure andiroba oil is mainly between 300 °C and 400 °C, while above 500 °C andiroba oil is entirely decomposed. The TGA curves have different behavior above and below the temperature of the oil degradation for andiroba oil/chitosan films. Below the temperature of oil degradation, the oil seems to

Table 3. Thermal analysis of chitosan, andiroba oil, and andiroba oil/chitosan film in an N2 atmosphere.Concentration (% w/w) % weight lossChitosan Andiroba oil 100 °C 200 °C 300 °C 400 °C 500 °C 600 °C 700 °C

1.0 0.0 12 17 43 59 63 66 690.1 14 22 44 59 66 71 750.5 10 15 34 54 74 76 781.0 9 13 30 54 77 80 82

2.0 0.0 12 19 44 59 63 66 670.1 14 21 44 59 65 69 740.5 10 17 38 52 68 72 751.0 10 15 32 48 73 76 79

Chitosan powder 4 6 32 50 56 60 62Andiroba oil 0 1 2 21 100 100 100

Figure 3. TGA thermograms of andiroba oil and chitosan powder.

Figure 4. TGA thermograms of 1% and 2% (w/w) chitosan films containing andiroba oil.

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Polímeros, 26(2), 168-175, 2016 173

stabilize the film, but when the temperature is higher than the temperature of the oil degradation, the oil appears to promote the opposite. If we compare the weight loss up to 400 °C (Figure 4 and Table 3), the increase of andiroba oil concentration decreases the weight loss values (the TGA curves are in a superior position for 0.5 and 1.0% of andiroba oil). However, when the temperature is above 400 °C, there is one inversion and the TGA curves become inferior for these higher concentrations of andiroba oil. We also observed in Figure 4 that the incorporation of andiroba oil in chitosan film tends to shift the thermal degradation zone to higher temperatures. Such change is attributed to an increase in thermal stability by the incorporation of andiroba oil, and they were also supported by the differential thermogravimetric analysis - DTG (data not shown). Regardless, Figure 4 shows that both chitosan concentrations presented the same behaviour when the andiroba oil concentration increases.

Pelissari et al.[33] verified that the addition of oregano essential oil to chitosan-starch films did not influence the thermal stability of these films; however, they observed an increase in residue percentage after the incorporation of the oregano essential oil.

3.5 Differential scanning calorimetry

Figures 5, 6 shows the DSC results for the chitosan powder, andiroba oil, and andiroba oil/chitosan films. The results obtained from the DSC include the temperatures, and their respective ΔH values are presented in Table 4.

The DSC heating curve for pure andiroba oil showed one endothermic peak at 42.57 °C and some exothermic peaks between 180 °C and 360 °C. Oils are one complex mixture of triacylglycerols (TAGs) acting also as a solvent for minority components, such as vitamins, pigments, phenolic compounds, phospholipids, free fatty acids, and mono- and diacylglycerols[37]. The five top TAGs present in Andiroba oil are, in descending order: Palmitic-Oleic-Oleic, Palmitic-Palmitic-Oleic, Palmitic-Oleic-Stearic, Oleic-Oleic-Oleic/Stearic-Oleic-Linoleic, and Palmitic-Linoleic-Oleic[8,37]. Thus, the first endothermic peak is probably associated with the fusion of free fatty acids. The exothermic peaks are probably associated with the decomposition of TAGs.

Matos[38] conducted a thermal study of pure fatty acids by DSC. He observed endothermic peaks for melting temperatures at 64.55 °C ± 0.67 and 71.36 °C ± 0.30, respectively, for the palmitic acid and stearic acid. Exothermic peaks of decomposition were observed at 235.24 °C ± 7.57, 241.87 °C ± 5.61, 238.79 °C ± 12.11, and 268.26 °C ± 18.16, respectively, for the palmitic, stearic, oleic, and linoleic acids. These fatty acids also showed a sequence of many exothermic peaks above 350 °C[38].

The exothermic peaks between 200 °C and 300 °C are probably due to the decomposition of the TAGs and fatty acids into smaller chains while the sequence of peaks above 350 °C can be attributed to the decomposition of these compounds into even smaller chains. The chitosan-andiroba oil film presented two peaks (T1 and T2) while the andiroba oil-free film showed also two peaks (T1 and T2) with a shoulder on the descending side. These shoulders could be attributed to the decomposition described above (T > 350 °C). The increase in andiroba oil concentration promotes a decrease

in the temperature T1, which is attributed to the evaporation of water associated with the chitosan and could also be attributed to the hydrophobic properties of the andiroba

Figure 5. DSC run curves for andiroba oil and chitosan powder.

Figure 6. DSC run curves for 1% and 2% (w/w) chitosan films containing andiroba oil.

Kimura, V. T., Miyasato, C. S., Genesi, B. P., Lopes, P. S., Yoshida, C. M. P., & Silva, C. F.

Polímeros, 26(2), 168-175, 2016174

oil. Such decrease in this temperature was, even more, prominent for 1% chitosan films. This fact can be explained by the hydrophobic character of andiroba oil and also by the reduction in enthalpy related to this evaporation (ΔH1).

The temperature T2 for the decomposition of chitosan is little affected by the addition of andiroba oil; however, the enthalpy (ΔH2) of the chitosan decomposition is reduced, which shows that andiroba oil reduces the heat generated by the decomposition of chitosan, thus confirming the increased thermal stability of the film that was observed in the TGA results.

4. Conclusions

Chitosan films containing andiroba oil were obtained with satisfactory properties for use as a dressing to heal wounds. The addition of andiroba oil significantly modified the mechanical and barrier properties of the films and promoted greater thermal stability for the films. About the barrier properties, it was found that all of the films exhibited properties which were compatible with commercial dressings, and the films with the highest concentration of chitosan showed best results. Concerning the mechanical properties, the Young’s modulus and tensile strength increased with the addition of andiroba oil; however, as the concentration of andiroba oil increases, the values for these properties approach the values observed for the film without andiroba oil. While elongation remained practically unchanged for the films with lower concentrations of chitosan (1.0%), for the films with higher concentrations of chitosan, the addition of andiroba oil promoted a decrease in elongation.

5. Acknowledgements

The authors wish to thank São Paulo Research Foundation (FAPESP) for the financial support (2010/17.721-4).

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Table 4. Temperatures and enthalpy measured by DSC.Concentration (% w/w) Temperature (°C) Enthalpy (J/g)

Chitosan Andiroba oil T1 T2 ΔH1 ΔH2

1.0

0.0 103.98 283.02 –288.99 233.480.1 83.60 288.00 –297.25 268.720.5 66.51 288.88 –210.18 181.101.0 61.29 289.43 –112.70 155.37

2.0

0.0 91.07 282.37 –340.13 266.230.1 63.36 287.69 –203.67 102.580.5 77.92 282.47 –259.22 116.441.0 69.55 283.72 –228.50 138.21

Chitosan powder 85.81 296.68 –248.53 198.83

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Polímeros, 26(2), 168-175, 2016 175

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35. Pranoto, Y., Rakshit, S. K., & Salokhe, V. M. (2005). Enhancing antimicrobial activity of chitosan films by incorporating garlic oil, potassium sorbate and nisin. LWT - Food Science and Technology, 38, 859-865. http://dx.doi.org/10.1016/j.lwt.2004.09.014.

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37. Saraiva, S. A., Cabral, E. C., Eberlin, M. N., & Catharino, R. R. (2009). Amazonian vegetable oils and fats: fast typification and quality control via triacylglycerol (TAG) profiles from dry matrix-assisted laser desorption/ionization time-of-flight (MALDI-TOF) mass spectrometry fingerprinting. Journal of Agricultural and Food Chemistry, 57(10), 4030-4034. http://dx.doi.org/10.1021/jf900043u. PMid:19358529.

38. Matos, F. C. (2012). Study of thermal decomposition of fatty acids through differential scanning calorimetry (Master’s dissertation). Universidade Estadual de Campinas, Campinas.

Received: Nov. 20, 2014 Revised: Dec. 24, 2015

Accepted: Feb. 15, 2016

http://dx.doi.org/10.1590/0104-1428.2151

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Preparação e caracterização de poliuretanos contendo diferentes quantidades de óleo de baru

Preparation and characterization of polyurethane with different quantities of baru oil

Elizabeth Luiza de Almeida1, Gilberto Alessandre Soares Goulart2, Salvador Claro Neto3, Gilberto Orivaldo Chierice3 e Adriano Buzutti de Siqueira1*

1Laboratório de Estudos de Materiais – LEMat, Universidade Federal de Mato Grosso – UFMT, Cuiabá, MT, Brasil

2Escola de Engenharia de Alimentos – EA, Universidade Federal de Goiás – UFG, Goiânia, GO, Brasil3Instituto de Química de São Carlos – IQSC, Universidade de São Paulo – USP, São Carlos, SP, Brasil

*buzutti7@hotmail.com

Resumo

O óleo de baru (OB) é uma matéria-prima renovável proveniente do cerrado brasileiro que possui elevado percentual de insaturação. O objetivo principal do trabalho foi o estudo das espumas de poliuretano com diferentes quantidade de óleo de baru saponificado. Os espectros de FTIR mostraram bandas representativas do poliol, óleo de baru e do pré-polímero. As curvas TG-DSC nas atmosferas de ar e N2 apresentaram perfis diferentes, devido a formação de produtos intermediários de oxidação dos materiais na atmosfera de ar. A quantidade de energia liberada pela DSC e as perdas de massa observadas pelas curvas TG indicaram saturação do óleo de baru na quantidade de 24,00% (m/m). Os gases liberados analisados por TG-DSC acoplado ao FTIR, foram: água, CO2, CO e etanol. Com o aumento do óleo de baru, ocorreu maior absorção de água pelo PU, na proporção de 29,00% (m/m) o índice de intumescimento foi de 91,10%, em 1440 minutos. Assim, pudemos observer a formação do poliuretano com características distintas devido a presença do OB.

Palavras-chave: poliuretano, óleo de baru, TG/DSC, FTIR.

Abstract

The baru oil (OB) is a renewable raw material from the brazilian cerrado, with high unsaturation percentage. The objective of this work was the study of polyurethane foam with different quantities of baru saponified oil. FTIR spectrum showed the bands of polyol, baru oil and prepolymer. The TG-DSC curves in N2 and air indicated different profiles, especially the formation of more stable intermediate products on the oxidation of these materials in air atmosphere. The quantity of heat obtained by DSC curves and the weight loss obtained by TG curve, indicates an OB saturation at 24.00% (w/w). The study by FTIR of the volatile products released by TG-DSC were water, CO2, CO and etanol. With the increase of OB, higher water absorption of PU, in 29.00% (w/w) of swelling index 91.10% for 1440 minutes. From these results, we observe the formation of polyurethane with different characteristics because of OB presence.

Keywords: polyurethane, baru oil, TG/DSC, FTIR.

1. Introdução

Os óleos vegetais, formados principalmente por triglicerídeos, fornecem uma excelente plataforma para síntese de materiais poliméricos. Devido a sua baixa toxicidade, baixo custo de produção e processamento, e por serem geralmente biodegradáveis. Vários tipos de óleos vegetais estão sendo utilizados para síntese de polímeros, tais como: soja, canola, girassol e linhaça, devido as suas excelentes propriedades[1,2].

O baru (Dipteryx alata Vog.), árvore da família Fabaceae, disseminada no Bioma Cerrado, faz parte do grupo das

espécies nativas usadas pela população regional como fonte de renda familiar[3]. Apresenta frutos do tipo drupa, ovóides, levemente achatados e de coloração marrom, com uma única semente (amêndoa) comestível e comercializada em empórios nos grandes centros, bastante apreciada pela população local. Se prolifera na região do Planalto Central, precisamente no norte de Minas Gerais, Goiás e centro de Mato Grosso, indo até a costa atlântica do Maranhão[4-6].

Em estudo realizado por Takemoto et al.[4], a semente de baru apresentou teores relativamente elevados de lipídios

Preparação e caracterização de poliuretanos contendo diferentes quantidades de óleo de baru

Polímeros, 26(2), 176-184, 2016 177

(38,2g/100g) de proteínas (23,9g/100g) e de calorias (502kcal/100g), além de fibras alimentares (13,4g/100g) e de minerais, como potássio (827mg/10g), fósforo (358mg/100g) e magnésio (178mg/100g). O óleo da semente revelou um elevado grau de insaturação (81,2%), conteúdo de α-tocoferol (5,0mg/100 g) e composição em ácidos graxos semelhantes a do óleo de amendoim, destacando-se os ácidos oléicos (50,4%) e linoléico (28,0%), este considerado essencial, o que favorece seu uso para fins alimentícios e como matéria-prima para as indústrias farmacêutica e oleoquímica.

Poliuretano (PU) é um dos materiais poliméricos mais versáteis, em relação aos seus métodos de processamento e suas propriedades mecânicas, podendo-se obter um plástico rígido cristalino, um elastômero flexível ou um gel viscoelástico[7,8]. Podendo ainda ser utilizado como espumas (flexíveis, rígidas e semi-rígidas), revestimentos, adesivos e fibras[9].

O objetivo do trabalho foi preparar e caracterizar os poliuretanos obtidos a partir da mistura do pré-polímero MDI (difenilmetano diisocianato), do poliol de óleo de mamona adicionando óleo de baru em diferentes proporções de massa (5,00%, 15,00%, 24,00% e 29,00%).

2. Materiais e Métodos

Todos os reagentes utilizados são de grau analítico, sendo o NaOH da marca Synth com grau de pureza 99,9% e etanol da marca Synth P.A. A obtenção do óleo de Baru bruto (OBB) ocorreu a partir de extração das amêndoas do baru, armazenadas a -18ºC, em hexano. O pré-polímero (PP) e o poliol utilizados na síntese foram desenvolvidos pelo Grupo de Química Analítica e Tecnologia de Polímeros-GQATP, USP de São Carlos-SP e produzidos pela indústria Cequil. O pré-polímero é baseado em difenilmetano diisocianato (MDI) e o poliol de óleo de mamona.

2.1 Preparação dos poliuretanos

O OBB foi saponificado em balão de fundo redondo, foi adicionado 120,0mL de água destilada em mistura contendo OBB(20,0g)/NaOH(4,00g)/CH3CH2OH(40,0mL). O balão de fundo redondo foi mantido por 2 horas colocado em rotaevaporador da LOGEN Scientific, modelo LSCS-1/52. A mistura foi transferida para béquer, sendo ajustado para valor de pH igual a 4 com adição de solução aquosa de ácido sulfúrico 30,0% (v/v)[10].

A mistura do óleo de baru saponificado (OBS) com o PP e o poliol foi realizado por agitação manual de 18,00g PP e 2,00g poliol em respectivos 1,00g (5,00%), 3,00g (15,00%), 5,00g (24,00%) e 6,00g (29,00%) de OBS.

2.2 Métodos de caracterização

Os espectros de absorção na região do infravermelho com Transformada de Fourier (FTIR) foram obtidos no Espectrofotômetro Perkin Elmer, modelo Spectrum 100, com resolução 4cm–1, na região compreendida entre 4.000 – 600cm–1, utilizando a técnica de reflectância total atenuada (ATR) com cristal de diamante/ZnSe.

As curvas TG-DSC foram obtidas no equipamento TG/DSC-1 da Mettler Toledo. O sistema foi calibrado

seguindo as especificações fornecidas pelo fabricante. As curvas foram obtidas em cadinho de α-Al2O3 (70 μL), com massa de amostra de aproximadamente 6 mg, razão de aquecimento de 20 °C min–1, atmosfera de ar seco e N2 com vazão de 60mL min–1 e intervalo de temperatura de 30-1000 °C.

As curvas TG-DSC e os espectros de infravermelho dos produtos gasosos foram obtidas no equipamento TGA/SDTA 851 da Mettler Toledo acoplado ao sistema de espectrometria infravermelha com transformada de Fourier, modelo Nicolet iS10 FTIR.

Para a determinação do percentual de intumescimento (Ii%), os PU’s foram cortados em pedaços de 1cm3 e mantidas em dessecador por 7 dias, os testes realizados em triplicata. Os PU’s foram submersos em água destilada, e em cada intervalo de tempo (1 min, 2 min, 5 min, 10 min, 15 min, 20 min, 30 min, 60 min e 1440 min), as peças eram cuidadosamente removidas, utilizando-se pinça, sendo retirado o excesso de água com papel toalha e, em seguida, as massas dos PU’s eram medidas. Este procedimento foi realizado de acordo com Soares[11] e Pereira et al.[12].

Foi construído um gráfico de intumescimento gravimétrico, utilizando-se a equação abaixo:

100%

Massa Final Massa Inicial xIiMassa Inicial

−= (1)

3. Resultados e Discussão

O óleo de baru foi submetido à saponificação para a liberação de hidroxilas nas cadeias, além de tornar o óleo miscível, este processo facilita a reação com o PP e o poliol na reação de polimerização. O aspecto visual dos PUs com OB indicaram tamanho de poros diferentes, indicando a influência da quantidade de OB no processo de polimerização.

3.1 Análise térmica do PP, Poliol, OBB e OBS

3.1.1 TG-DSC do OBB

As curvas TG-DSC do OBB mostram a decomposição térmica em 3 (atmosfera N2) e 4 (atmosfera de ar) etapas consecutivas, ver Figuras 1 e 2 e Tabelas 1 e 2. A primeira etapa de decomposição térmica são equivalentes a saída de água adsorvida do ambiente, sendo observado perda de massa inferior a 1%, indicando ser material não higroscópico. A segunda etapa de decomposição térmica são referentes a saída de compostos orgânicos voláteis de baixa massa molecular e provavelmente a evaporação do ácido gadoléico, presentes no OBB[4].

Em atmosfera de ar foram observados eventos exotérmicos na terceira etapa de decomposição térmica referentes a oxidação de compostos presentes. A terceira etapa de decomposição térmica, obervada em atmosfera de N2, foi devido a evaporação de compostos com massas molares maiores (C16:0, C18:0, C18:1, C18:2, C20:0, C22:0), ácido lignocérico e compostos de α e γ-tocoferol. A formação de produtos intermediários oxidados em ar, diferente da evaporação, inferiu maior temperatura final (Tf = 647 °C) na decomposição térmica do OBB. Os resíduos estáveis formados até 1000ºC, 1,15% (ar) e 1,20% (N2), são referentes as impurezas (material inorgânico) dos materiais.

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Figura 1. TG-DSC das matérias-primas dos poliuretanos em atmosfera de ar.

Tabela 1. Dados termoanaliticos das matérias-primas dos poliuretanos em atmosfera de ar seco.

AmostraPrimeira etapa Segunda etapa Terceira etapa Quarta etapa

T °C Δm% Pico/ºC T °C Δm% Pico/ °C T °C Δm% Pico/ °C T °C Δm% Pico/ °COB 67-140 0,69 - 140-250 4,37 225(↑) 250-478 88,53 338(↑) 478-647 5,14 616(↑)

OBS 36-85 6,88 52(↓) 85-138 16,5290, 99,

117, 123 e 128(↓)

138-502 68,20369,

386, 423, 474(↑)

502-650 4,57524,

543, 595, 634(↑)

Poliol 47-174 3,93 105(↓) 174-489 90,35 298, 397, 467(↓) 489-685 4,18 635(↓) - - -

PP 135-339 58,32 306(↑) 339-784 38,62 628, 690(↓) - - - - - -

Tabela 2. Dados termoanaliticos das matérias-primas dos poliuretanos em atmosfera de N2.

AmostraPrimeira etapa Segunda etapa Terceira etapa Quarta etapa

T °C Δm% Pico/ °C T °C Δm% Pico/ °C T °C Δm% Pico/ °C T °C Δm% Pico/ °COB 70-144 0,84 - 144-250 3,34 185(↑) 250-492 94,26 433(↓) - - -

OBS 51-81 6,10 38,53(↓) 81-143 15,92 108,119, 120 (↓) 143-516 72,35 301,394,

459(↓) 535-800 1,97 726(↓)

Poliol 52-181 3,85 - 181-497 93,72 381,446(↓) - - - - - -

PP 139-354 56,90 278(↓) 354-1000 19,01 542(↓) - - - - - -

3.1.2 TG-DSC do OBS

As curvas TG-DSC do OBS em atmosfera de ar e N2 apresentaram perfis similares até 350 °C, a distinção das curvas depois desta temperatura é decorrente da maior

estabilidade térmica dos produtos intermediários oxidados formados em atmosfera de ar. As curvas TG do OBS até 138 °C (ar) e 143 °C (N2) (Δmar=33,41% e ΔmN2= 31,85%), estão relacionadas desidratação do óleo (água de hidratação

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e constituição) provindas do processo de saponificação do OBB. Os resíduos formados em 650 °C (ar) e em 750 °C (N2) foram, respectivamente, 3,33%(m/m) e 3,38%(m/m), estes resíduos são maiores do que os observados em OBB, devido a inserção de NaOH.

3.1.3 TG-DSC do Poliol

Apesar da similaridade nas curvas TG-DSC do poliol do óleo de mamona com o OBB, pode ser notado no terceiro evento de decomposição térmica, que a quantidade de calor liberado, obtido pela curva DSC, no OBB (ΔH=1,01kJ g–1) foi significativamente maior que a observada na curva DSC do poliol (ΔH= 0,40kJ g–1), indicando característica intrínseca de cada matéria-prima estudada, e portanto podendo utilizar este parâmetro para distinguir aos dois óleos.

3.1.4 TG-DSC do Pré-polímero

A decomposição térmica do PP anidro, em ambas as atmosferas, ocorreram em duas etapas, com indicação de patamar entre 300-400 °C. Em atmosfera de N2, houve a formação de resíduo carbonáceo (Δmtotal=76,07%), e em atmosfera de ar ocorreu a degradação oxidativa da matéria orgânica, com perda de massa igual a 97,18% do pré-polímero. O resíduo estável até 1000 °C, em atmosfera de ar, foi decorrente da presença de impureza de compostos inorgânicos do pré-polímero, podendo ser estimado a sua pureza em no máximo 97% (m/m).

3.2 Análise Térmica do PU, PU-OBS 5,00%, PU-OBS 15,00%, PU-OBS 24,00% e PU-OBS 29,00%

3.2.1 TG-DSC do PU

As curvas TG-DSC da PU em atmosfera de ar seco e N2 apresentaram perfis distintos, devido as oxidações decorrentes da atmosfera em ar seco, ver Figuras 3 e 4 e as Tabelas 3 e 4. A primeira perda de massa é relativo a saída de água de hidratação da PU, indicando que o material não é higroscópico. Teste qualitativo (visualização do resíduo) e a menor perda de massa (∆m-ar = 98,90% e ∆m-N2 = 90,50%), sugere a formação de resíduo carbonáceo em atmosfera de N2. A massa residual em atmosfera de ar seco retrata a impureza da PU, provavelmente em função do material inorgânico presente nas matérias primas.

3.2.2 TG-DSC do PU-OBS 5,00%, PU-OBS 15,00%, PU-OBS 24,00% e PU-OBS 29,00%

Os eventos observados nas curvas TG-DSC das quatro amostras de PUs contendo OBS (5,00%, 15,00%, 24,00% e 29,00%), em atmosfera de ar seco e N2, são mostrados nas Figuras 5 e 6 e Tabelas 3 e 4, sendo que os perfis das curvas apresentaram faixas de temperatura semelhantes. Apesar da estabilidade térmica da PU contendo OBS ser maior que a PU sem OBS, foi verificado que a estabilidade térmica das PU-OBS diminuem à medida que a proporção em massa de OBS aumenta.

As curvas TG-DSC da espuma de PU sem OBS, apresentaram 2 picos exotérmicos em 320 °C e 560 °C, em atmosfera de ar seco, com perda de massa igual a 98,15% (Figura 3). As perdas de massa das PUs com OBS foram 96,16% (5,00% OBS), 96,83% (15,00% OBS), 94,35%

(24,00% OBS) e 94,46 (29,00% OBS), sendo observados nas curvas DSCs 3 picos exotérmicos em aproximadamente 350 °C, 580 °C e 680°C, ocasionados pela decomposição oxidativa dos materiais. Os picos exotérmicos em torno de 680 °C, observados nas curvas TG-DSC dos PUs-OBS, são relativos a decomposição oxidativa dos OBS nas PUs. A observação dos picos exotérmicos em 700 °C (5,00% OBS), 681 °C (15,00% OBS), 680 °C (24,00% OBS) e 678 °C (29,00% OBS), sugere a interação dos OBS durante o processo de polimerização dos PUs. A quantidade de energia

Figura 2. TG-DSC das matérias-primas dos poliuretanos em atmosfera de N2.

Figura 3. TG-DSC do poliuretano em atmosfera de ar.

Figura 4. TG-DSC do poliuretano em atmosfera de N2.

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calculado através das curvas DSC, (ΔH5%OBS=92 J g–1; ΔH15%OBS=505 J g–1; ΔH24%OBS= 942 J g–1; ΔH5%OBS= 850 J g–1), assim como as perdas de massa observadas na quarta etapa de decomposição térmica da curva TG, indicam saturação do OBS nas reações de polimerização, quando utilizado 24% (m/m).

Nas amostras em atmosfera de N2 (Figura 6) foi observado um pico endotérmico em 360 °C ocasionado pela decomposição térmica. O evento exotérmico com grande liberação de energia entre 600 °C e 820 °C, ocorreu devido a formação dos gases liberados na decomposição térmica, gerando reações oxidativas no sistema. Em função das

Tabela 3. Dados termoanalíticos dos PU’s com diferentes concentrações de OBS em atmosfera de ar seco.

AmostraPrimeiro evento Segundo evento Terceiro evento Quarto evento

T °C Δm% T °C Δm% T °C Δm% T °C Δm%PU 35-179 0,75 179-356 24,45 356-481 14,53 481-695 59,17PU-OB-5% 33-108 1,18 250-375 24,19 375-681 69,30 681-700 2,67PU-OB-15% 36-92 0,73 200-364 23,37 364-650 65,50 650-680 7,23PU-OB-24% 33-84 0,74 190-359 24,06 359-643 59,98 643-680 10,31PU-OB-29% 33-92 0,86 180-363 23,68 363-646 60,90 646-693 9,88

Tabela 4. Dados termoanalíticos dos PU’s com diferentes concentrações de OBS em atmosfera de N2.

AmostraPrimeiro evento Segundo evento Terceiro evento Quarto evento

T °C Δm% T °C Δm% T °C Δm% T °C Δm%PU 34-173 0,65 173-372 29,27 372-540 57,97 540-1000 2,58PU-OB-5% 33-108 1,04 209-378 38,30 378-553 39,13 553-1000 4,70PU-OB-15% 34-112 0,80 188-381 38,62 381-551 39,29 551-1000 6,09PU-OB-24% 33-84 0,52 154-366 33,87 366-541 43,21 541-1000 9,00PU-OB-29% 29-88 0,55 154-366 33,43 366-542 42,27 542-1000 10,00

Figura 5. TG-DSC dos poliuretanos em atmosfera de ar.

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reações de oxidação observadas nas curvas TG-DSC, não houve formação de resíduo carbonáceo.

3.3 Espectroscopia de absorção na região do infravermelho

As bandas observadas no espectro FTIR do OBB e OBS (Figura 7), foi verificado algumas semelhanças nas vibrações moleculares, com exceção da banda em 3380cm–1, correspondente as interações de hidrogênio das hidroxilas, relativas ao processo de saponificação. Em comparação com os resultados de Drummond 2008, os picos que aparecem em 2922cm–1 e 2853cm–1 correspondem aos estiramentos da ligação C-H, atribuídos principalmente aos componentes lipídicos do óleo. Na região de 1744cm–1 aparece um estiramento atribuído à ligação C=O, característico do grupo carbonilico do ácido graxo. A banda em 1160cm–1 e 721cm–1 são atribuídos, respectivamente a ligação C-O e C-H dos ácidos graxos. Na Figura 8 está representado o espectro do poliol, onde pode ser observado as principais bandas em:

I) 3364cm–1, atribuida a presença de hidroxila no óleo;

II) 1743cm–1, referente ao grupo C=O;

III) 2925cm–1 e 2845cm–1, atribuídos, respectivamente, aos estiramentos simétrico e assimétrico do grupo –CH2

[2,4-9,11-14].

Figura 6. TG-DSC dos poliuretanos em atmosfera de N2.

O espectro de FTIR do pré-polímero (Figura 8), o mesmo material utilizado por Trovati et al., 2010 na caracterização do poliuretano de óleo de mamona, apresentou perfil semelhante ao encontrado neste trabalho, com banda característica do grupo isocianato em 2241cm–1, e bandas indicando grupos uretanos polimerizados em 1718cm–1, 1608cm–1, 1577cm–1 e 1520cm–1, bandas típicas de estiramentos C=O e ligações N-H.

As principais bandas características de poliuretanos são as indicativas pela presença dos ésteres (C=O) do poliol, que estão em 1728cm-1 e 1124cm–1 e ainda em 814cm–1 e 1598cm–1 referentes aos isocianatos da estrutura do uretano[15,16]. A diminuição da intensidade da banda em 2250cm–1 e a formação da banda característica de grupos uretanos polimerizados evidenciaram a formação do poliuretano, tais como bandas intensas na região de 3421-3447 e 1721cm–1, referentes às ligações uretano, demonstrando consistência nas interações NCO-OH[17,18].

Para todos os PUs-OBS (Figura 8) foi observado a banda em 1738 cm–1, característico do poliol referente ao grupo carbonila, sendo que ocorreu diminuição em sua intensidade relativa, à medida que a proporção de OBS aumenta, nas proporções de 15,00%, 24,00% e 29,00%, indicando a inserção do OBS no PU. As bandas observadas em 3320cm–1 são referentes aos grupos hidroxilas presentes no poliol e no OBS.

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Figura 7. Espectros FTIR do (a) OBB e (b) OBS.

Figura 8. Espectros FTIR do: Polol, PP, PU-OBS 5%, PU-OBS 15%, PU-OBS 24% e PU-OBS 29%.

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Uma análise comparativa dos espectros das matérias primas, do PU e so PU-OBS sugerem ocorrência de polimerização com o OBS. Esta sugestão é devido ao deslocamento das bandas do espectro FTIR de 3354 cm–1 (OBS e poliol) para 3351cm–1 (PU) e 3333cm–1 (PUs-OBS-5%/15%/24%/29%), referentes aos estiramentos das hidroxilas. Os valores relativos de transmitância das PUs sugerem pouca alteração na quantidade de hidroxilas presentes na PU com as PUs-OBS, sendo que o valor observado no PU-OBS 29,00% é similar ao observado no PU.

3.4 Análise por FTIR dos gases liberados, acoplado ao TG-DSC

Os produtos gasosos detectados por FTIR na decomposição térmica dos poliuretanos, com ou sem OBS, foram CO2, H2O, etanol e CO, em ambas as atmosferas, ver Figura 9, com exceção do PU-OB-5% que apresentou liberação somente de CO2, característica dos grupos isocianato, entre 50 °C até 370 °C. Nas análises dos gases do PU, PU-OBS 15%, PU-OBS 24%, e PU-OBS 29%, ocorreu a liberação de CO2 na faixa de temperatura entre 30 °C até 410 °C, H2O com formação de uma banda em 3500cm-1, indicativo de

hidroxila, banda em 1500cm–1, típico de estiramento C=O e ligações N-H dos uretanos.

Na atmosfera de ar seco, todos os poliuretanos mostraram comportamento semelhante na liberação de produtos gasosos, com exceção do PU-OB-5% que não apresentou liberação de etanol na faixa de 230-300 °C, os demais tratamentos apresentaram além do etanol, CO2 na faixa de temperatura entre 300-700 °C. A saída de H2O pode ser observada em dois momentos, no início do processo de degradação entre 100-370 °C e nos momentos finais de 550-590 °C. A saída de CO pode ser observado entre 390-630 °C.

3.5 Indice de intumescimento

A capacidade de adsorção de água nos PUs, com e sem OBS, foi avaliado pelo índice de intumescimento, ver Tabela 5. Somente o sistema PU-OBS 15,00% apresentou ponto máximo de absorção de água aos 20 minutos (42,80%), nos outros sistemas estudados houve aumento no percentual do índice de intumescimento com o aumento da massa do OBS, ao final das 24 horas de imersão. O aumento no índice de intumescimento pode ser explicado pelo aumento da hidrofilicidade do sistema, pois a difusão de água é

Figura 9. Espectros FTIR dos gases liberados do PU-OB 29%, representativo a todos os PU-OBS.

Tabela 5. Índice de Intumescimento dos poliuretanos com diferentes concentrações do óleo de baru.

Tempo (min)Índice de Intumescimento %

PU-OB-5,00% PU-OB-15,00% PU-OB-24,00% PU-OB-29,00%1 11,2 9,5 41,0 39,93 15,8 15,7 45,3 23,55 23,7 14,6 38,6 39,77 27,9 20,4 30,4 41,310 31,9 18,9 34,3 28,715 32,9 11,8 44,8 22,720 25,7 42,8 47,9 24,330 25,1 32,2 46,5 21,260 25,1 32,9 45,9 21,7

1440 34,9 25,4 50,4 91,1

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determinada pela microestrutura do material e a afinidade dos componentes poliméricos pela água. .

Como o poliol favorece naturalmente a solubilidade em água devido a presença da hidroxila no carbono 12 da cadeia lipofílica, a adição do OBS no PU faz aumentar a absorção de água pelo poliuretano.

4. Conclusão

Através das curvas termoanalíticas, termogravimetria-calorimetria exploratória diferencial simultâneo (TG-DSC) podemos estudar as relações proporcionais das espumas formadas, estudar a estabilidade térmica e as etapas de decomposição térmica dos mesmos, assim como identificar a presença de impurezas nos materiais.

Apesar das curvas TG-DSC indicarem similaridades do OBB com o poliol, houve necessidade de saponificar o OBB para que o mesmo pudesse reagir na polimerização.

Os espectros de FTIR puderam confirmar a polimerização do PU, mesmo com a inserção de OBS, seugerindo inclusive grande similaridade nas propriedades dos mesmos.

O índice de intumescimento mostrou que o OBS aumenta a absorção de água pelo PU.

Tendo em vista a afinidade da PU-OBS por água, e sendo esta associação ocasionada pelas interações de hidrogênio água-OB, portanto, é provável que o sistema PU-OBS venha a interagir com metanol e etanol. Caso ocorra adsorção metanol e etanol, de forma seletiva, este sistema poderá ser utilizado na purificação do biodiesel.

5. Agradecimentos

A FAPEMAT e CNPq pelo suporte financeiro e Bolsa de mestrado concedido pela CAPES. A Engenheira de Alimentos Tatiane R. Silva por auxiliar na preparação do óleo de Baru bruto, utilizado neste trabalho.

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13. Drummond, A. L. (2008). Compósitos poliméricos obtidos a partir do óleo de Baru: síntese e caracterização (Dissertação de Mestrado). Instituto de Química, Universidade de Brasília, Brasília.

14. Trovati, G., Sanches, E. A., Claro, S. No., Mascarenhas, Y. P., & Chierice, G. O. (2010). Characterization of polyurethane resins by FTIR, TGA, and XRD. Journal of Applied Polymer Science, 115(1), 263-268. http://dx.doi.org/10.1002/app.31096.

15. Cangemi, J. M., Claro, S. No., Chierice, G. O., & Santos, A. M. (2006). Study of the biodegradation of a polymer derived from castor oil by scanning electron microscopy, thermogravimetry and infrared spectroscopy. Polímeros: Ciência e Tecnologia, 16(2), 129-135. http://dx.doi.org/10.1590/S0104-14282006000200013.

16. Pellizzi, E., Derieux, A. L., Lavédrine, B., & Cheradame, H. (2013). Degradation of polyurethane ester foam artifacts: chemical properties, mechanical properties and comparison between accelerated and natural degradation. Polymer Degradation and Stability, 2013, 1-7. http://dx.doi.org/10.1016/j.polymdegradstab.2013.12.018.

17. Ismail, E. A., Motawie, A. M., & Sadek, E. M. (2011). Synthesis and characterization of polyurethane coatings based on soybean oil–polyester polyols. Egyptian Journal of Petroleum, 20(2), 1-8. http://dx.doi.org/10.1016/j.ejpe.2011.06.009

18. Paiva, G. M. S., Borges, F. D. M., Batista, N. C., Lima, S. G. D., & Matos, J. M. E. D. (2010). Síntese de poliuretana obtida do óleo de mamona (ricinus communis) na ausência de solventes. In Anais do 62º Congresso Brasileiro para o Progresso da Ciência (pp. 1). Natal: SBPC.

Enviado: Jun. 03, 2015 Revisado: Ago. 03, 2015

Aceito: Set. 17, 2015

http://dx.doi.org/10.1590/0104-1428.1905

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Polímeros, 26(2), 185-191, 2016 185

Caracterização de pinos da blenda poli(L-co-D,L ácido láctico)/poli(caprolactona triol) (PLDLA/PCL-T) e análise das

propriedade mecânicas dos pinos durante degradação in vitro

Characterization of pin blends Poly (L-co-D,L lactic acid) /Poly(caprolactone diol) (PLDLA/PCL-T) and analysis of mechanical property of pins during in vitro degradation

Marcia Adriana Tomaz Duarte1,2, Adriana Cristina Motta1 e Eliana Aparecida de Rezende Duek1,3

1Departamento de Ciências Fisiológicas, Centro de Ciências Médicas e Biológicas, Pontifícia Universidade Católica de São Paulo – PUC-SP, Sorocaba, SP, Brasil

2Mestrado Profissional em Engenharia Mecânica, Centro Universitário SOCIESC, Santa Catarina, SC, Brasil

3Departamento de Engenharia de Materiais, Faculdade de Engenharia Mecânica, Universidade Estadual de Campinas – UNICAMP, Campinas, SP, Brasil

*eliduek@fem.unicamp.br

Resumo

Os dispositivos de fixação óssea, metálicos convencionais, usados em cirurgia crâniomaxilofacial têm apresentado alguns problemas, tais como, corrosão, inflamação e infecção, além de neoformação de estrutura óssea mecanicamente inferior devido à atrofia gerada pela diferença de módulo elástico entre metal e osso, razões que têm levado ao aumento do interesse por dispositivos poliméricos bioarreabsorvíveis. Os polímeros biorreabsorvíveis mais utilizados nesta aplicação pertencem à família dos poli (α-hidroxi ácidos), que têm como característica degradarem por hidrólise de suas ligações ésteres, tal como copolímero poli (L-ácido láctico-co-D, L ácido láctico), PLDLA. Neste trabalho foram investigados alguns efeitos da adição de poli (caprolactona triol), PCL-T sobre PLDLA. Foram preparados pinos por fusão de blendas nas seguintes composições 100/0, 90/10, 70/30 and 50/50 (m/m), PLDLA/PCL-T. Os pinos foram caracterizados por diferentes técnicas (DSC, MEV e ensaio mecânico). A degradação in vitro dos pinos foi investigada, sendo observado que a adição de PCL-T no PLDLA modificou suas propriedades mecânicas e morfológicas. Tais mudanças podem apresentar potencial para outras aplicações do material, onde a questão da flexibilidade se faça necessária.

Palavras-chaves: PLDLA/PCL-T, caracterização, propriedades mecânicas, in vitro.

Abstract

The bone fixation devices, conventional metallic cranium used in surgery have presented some problems, such as corrosion, inflammation and infection, and a lower mechanically newly formed bone structure due to the atrophy caused by the difference in stiffness between metal and bone. These reasons have led to increased interest in bioreabsorbable polymeric devices. The most bioresorbable polymers used in this application belong to the family of poly (α-hydroxy acids), which are characterized degrade by hydrolysis of its ester linkages, such as copolymer poly (L-lactic acid-co-D, L lactic acid) (PLDLA). In this work was investigated some effects of the addition poly (caprolactone triol), PCL-T in the PLDLA. Pins were prepared by melting in the 100/0, 90/10, 70/30 and 50/50 (w/w) compositions, PLDLA/PCL-T. The pins were characterized by different methods (DSC, SEM and mechanical test). The in vitro pins degradation was investigated. It was observed that the addition of PCL-T of PLDLA modifies the mechanical properties, morphological, we conclude that PLDLA/PCL-T can be a potential material for various applications, where the question of flexibility to make necessary.

Keywords: PLDLA/PCL-T, characterization, mechanical properties, in vitro.

Duarte, M. A. T., Motta, A. C., & Duek, E. A. R.

Polímeros, 26(2), 185-191, 2016186

1. Introdução

Os defeitos ósseos podem ser resultado de anormalidades congênitas, traumas ou doenças. As aplicações de dispositivos de fixação óssea interna a partir de biomateriais, utilizados nestas situações, revolucionaram o campo de consolidação de fraturas ósseas. Isto se deve às novas tecnologias disponíveis, resultantes do desenvolvimento de materiais que geram produtos que atendem às solicitações para uso biomédico.

A seleção do material é importante e deve garantir biocompatibilidade e resistência à corrosão durante o tempo necessário para recuperação do osso[1] e embora os materiais metálicos permitam bons resultados clínicos através de uma fixação estável do osso, seu emprego esta associado a algumas limitações e riscos, como a possibilidade de re-fratura ao longo do osso no local onde estava presente o dispositivo[2,3] em função da grande diferença de rigidez entre o osso (E = 1-30 GPa) e os metais (E = 100-200 GPa), ou ainda a possibilidade de reações inflamatórias por parte do tecido circunjacente ao implante metálico[4], causadas pela presença do metal por um longo período no corpo.

Dessa forma, os dispositivos metálicos, em inúmeras situações, devem ser removidos, por meio de uma 2o intervenção cirúrgica corroborando com a ideia de que não representam portanto os biomateriais ideais para aplicações ortopédicas, em função de reações alérgicas e liberação de íons metálicos ao redor do tecido implantado[5].

Neste contexto os poliméricos biorreabsorvíveis tem sido uma classe de materiais atrativa nas últimas décadas, sendo utilizados e experimentados em vários aspectos das cirurgias ortopédicas, incluindo fixações de fraturas, reposição óssea, reparo da cartilagem, reparo do menisco e fixação de ligamentos[6,7]. Esses materiais biorreabsorvíveis têm sido utilizados na forma de parafusos, pinos e placas nas aplicações ortopédicas e cirurgias orais em humanos e animais[8].

O implante de materiais poliméricos biorreabsorvíveis suporta uma elevada fração de carga transmitida ao osso durante os primeiros estágios do processo de recuperação. A redução gradual da resistência do implante, devido à sua reabsorção, transfere uma porcentagem crescente da carga ao osso em recuperação. Como resultado, o local da fratura recuperado desenvolve uma resistência comparável àquela original do osso[9].

Os poli (α-hidróxi ácidos) são considerados uma das famílias de polímeros mais promissoras na área dos bioreabsorvíveis. A grande vantagem desses polímeros está na sua forma de degradação que ocorre por hidrólise de suas ligações ésteres, sendo os produtos gerados completamente absorvidos pelo organismo. Dentre esses polímeros destaca-se o poli (L-ácido láctico-co-D, L ácido láctico) (PLDLA), em função do adequado tempo de degradação que apresenta, compatível ao requerido na fixação de fraturas ósseas.

O estudo tem caracterizado a retenção de resistência durante a degradação do copolímero PLDLA (70/30), que mantem 90% da resistência após 6 meses, 70% da resistência após 9 meses, 50% da resistência após 12 meses e nenhuma resistência após 18 meses[10]. A tecnica de auto-reforço inventada e patenteada por Törmälä e colaboradores[11], onde tanto a matriz quanto o elemento reforçante são do mesmo

polímero permite a obtenção de dispositivos poliméricos maleáveis e resistentes, sendo diversos os estudos de auto- reforço ao PLDLA.

Uma grande vantagem dessa técnica é a possibilidade da manufatura de dispositivos com tamanhos menores que podem ser flexionadas a frio. Comercialmente os copolímeros auto-reforçados PLDLA (70/30) (placas e parafusos) são bastante empregados (Biosorb FX, Bionximplants LTD, Tampere, Finlândia). Entretanto a técnica exige que sejam realizadas diversas etapas de processamento, o que representa uma dificuldade na obtenção do produto final.

Blendas de PLDLA com PCL-triol têm sido alvo de estudo na busca por aliar as características mecânicas do PLDLA com a flexibilidade do PCL triol. O PCL-triol age como um plastificante, enfraquecendo as forças intermoleculares entre as cadeias poliméricas[12], com objetivo de melhorar a processabilidade e aumentar a flexibilidade. Essa característica que o PCL-T atribui ao PLDLA em relação à flexibilidade pode ser interessante em diversas aplicações, como por exemplo, na engenharia tecidual. Espósito[13] realizou um estudo no qual a blenda PLDLA/PCL-T servia como scaffolds para crescimento celular visando aplicação em regiões como menisco.

Com isto, o objetivo deste estudo foi obter materiais que atendam a solicitação mecânica para fixação óssea sem a necessidade de auto-reforço, preparando blendas de PLDLA com PCL-triol em diversas proporções, sendo avaliadas as características da blenda. Foi acompanhado, o perfil da degradação in vitro.

2. Parte Experimental

2.1 Materiais

O poli (L-co-DL, ácido lático) (70:30) PLDLA, foi sintetizado nos laboratório da PUC/ Sorocaba, com massa molar média de 265000g.mol–1 de acordo com Motta[14]. A poli (caprolactona triol), (PCL-T), cujo nome químico é 2-oxepanona, foi fornecida pela Solvay (CAPA 3091), com massa molar média de 900g.mol–1.

2.2 Preparação dos pinos por fusão

Foram preparadas blendas de PLDLA/PCL-T por fusão, obtendo-se pinos com 2 mm de diâmetro, nas proporções 100/0,90/10,70/30,50/50% (m/m), utilizando uma mini-injetora Mini Max Molder modelo LMM – 2017. (Figura 1)

As amostras foram colocadas na mini injetora, e aquecidas a 200 °C por 1 minuto sem cisalhamento, seguido de 1,5 min com cisalhamento (velocidade de cisalhamento constante, de 4 rpm). O molde foi envolto por uma camisa para mantê-lo aquecido a 110 °C. Após esse procedimento a mistura foi injetada no molde, deixado esfriar a temperatura ambiente. Depois de retirado do molde o pino foi armazenado no dessecador.

2.3 Caracterização do PLDLA/PCL-T obtido por fusão

Calorimetria exploratória diferencial (DSC): As análises de DSC foram realizadas em um equipamento modelo 2920 da TA Instruments. Amostras pesando aproximadamente 7-10mg foram aquecidas de 25 °C a 200 °C a uma taxa

Polímeros, 26(2), 185-191, 2016 187

Caracterização de pinos da blenda poli(L-co-D,L ácido láctico)/poli(caprolactona triol) (PLDLA/PCL-T) e análise das propriedade mecânicas dos pinos durante degradação in vitro

de aquecimento de 10 °C.min–1 (primeiro aquecimento), e mantidas a esta temperatura por 5 min. Subsequentemente foram resfriadas a –100 °C a uma taxa de 10 °C.min–1 e mantidas nessa temperatura por 5 min. Após, foram aquecidas novamente ate 200 °C a uma taxa de 10 °C/min sob atmosfera de nitrogênio.

Microscopia eletrônica de varredura (MEV): As amostras das superfícies superiores e das fraturas (obtidas em nitrogênio líquido) dos pinos foram metalizadas com ouro (Sputter Coater BAL-TEC SCD 050) e analisadas em um microscópio eletrônico de varredura (JEOL JXA 860) operado a 10kV.

Ensaio mecânico de Flexão: Os pinos de PLDLA/PCL-T nas composições (100/0), (90/10), (70/30), (50/50)% (m/m) com cerca de 1,94 (± 0,45) mm de diâmetro foram submetidos a ensaios de flexão em equipamento servo-hidráulico da Testar II (modelo 810 MTS), célula de carga com capacidade de 1kN e veloc.50 mm.min–1 com vão de 30mm. Foram ensaiados 5 corpos de prova de cada amostra nas mesmas condições de umidade (50%) e temperatura (25 °C ± 2).

2.4 Degradação in vitro – análise macroscópica e comportamento mecânico dos pinos PLDLA/PCL-T

Os pinos de PLDLA/PCL-T nas composições 100/0, 90/10, 70/30 e 50/50 foram colocados em tubos de ensaios com tampa rosqueada, previamente esterilizados contendo solução tampão fosfato (PBS) pH 7,4 a 37 ± 1 °C, sendo retirados após 4, 8, 12 e 24 semanas, lavados com água destilada e secos sob vácuo, durante 48 horas. Depois

de secas as amostras foram caracterizadas pelas técnicas descritas a seguir:

Ensaio mecânico de Flexão: Os pinos de PLDLA/PCL-T nas composições (100/0), (90/10), (70/30), (50/50)% (m/m) com cerca de 1,94 (± 0,45) mm de diâmetro foram submetidas a ensaios de flexão em um equipamento servohidráulico da Testar II, modelo 810, do fabricante MTS, célula de carga com capacidade de 1kN e velocidade de 50mm.min–1 com vão de 30mm. As amostras foram submetidas a 5 ensaios nas mesmas condições de umidade (50%) e temperatura (25 °C ± 2).

Microscopia eletrônica de varredura (MEV): Foi acompanhado por meio da analises de MEV o processo de degradação dos pinos.

Avaliação Macroscópica: Foi realizada a avaliação macroscópica das amostras em função do tempo de imersão na solução tampão fosfato para as diferentes composições de PLDLA/PCL-T.

3. Resultados

3.1 Caracterização das composições PLDLA-PCL-triol

Calorimetria exploratória diferencial (DSC): Os resultados referentes ao DSC do sistema PLDLA/PCL-T nas composições 100/0, 90/10, 70/30 e 50/50 estão sumarizados na Tabela 1.

As análises de DSC conforme Tabela 1 mostram as temperatura de transição vítrea (Tg),para os polímeros porém, apenas o PCL-T apresenta temperatura de fusão (Tm),

Figura 1. Esquema de funcionamento da mini injetora Mini Max Moulder[15].

Tabela 1. Resultados da analise de DSC para as várias composições PLDLA-PCL-triol.

Composição(PLDLA/PCL-T)

AquecimentoTg

PCL-T(°C)

Tc

PCL-T(°C)

ΔHc

PCL-T(J/g)

Tm

PCL-T(°C)

ΔHm

PCL-T(J/g)

TgPLDLA

(°C)100/0 1° - - - - - 57100/0 2° - - - - - 5790/10 1° - - - - - 4290/10 2° –67 - - - - 4870/30 1° - - - - - 4070/30 2° –73 –32 2 1 1 4850/50 1° - - - - - 4150/50 2° –67 –32 2 3,4 0,5 46

Duarte, M. A. T., Motta, A. C., & Duek, E. A. R.

Polímeros, 26(2), 185-191, 2016188

caracterizando o PLDLA como polímero amorfo. O pino de PLDLA puro mostrou uma Tg de 57 °C no primeiro e segundo aquecimento conforme Figura 2a, b.

A adição de 10% de PCL-T diminuiu a Tg do PLDLA, de 57 °C para 48 °C, sendo que as demais composições de PCL-T mantiveram esse mesmo patamar de diminuição, independente da quantidade de PCL-T ser maior.

Concentração do PCL-T maiores que 10%, não modificaram significativamente o comportamento térmico do PLDLA, o que pode caracterizar como sendo esse um valor limite de concentração do PCL-T, acima do qual não influencia nas propriedades térmicas desse material conforme observado na Tabela 1.

3.2 MEV – Microscopia eletrônica de varredura (Pinos)

A Figura 3 mostra as micrografias obtidas por MEV das topografias das superfícies e das superfícies de fratura do PLDLA/PCL-T em várias composições obtido por fusão. O copolímero PLDLA em forma de pino apresenta uma topografia de superfície densa, lisa apresentando algumas incrustações (Figura 3a) provenientes do processo de fusão do material. Na análise microscópica da superfície de fratura é possível observar um material compacto, denso com ausência de poros (Figura 3b).

A composição 90/10 apresentou uma morfologia de superfície similar ao copolímero puro, conforme demonstrado na Figura 3c. A observação da fratura nesta composição

evidenciou a presença de poros, Figura 3d. Os poros estavam distribuídos homogeneamente e apresentavam diâmetros de 0,75-1μm, em formato esférico.

Os pinos com composição 70/30 (Figura 3e, f) e 50/50 (Figura 3g, h) tem topografia de superfície e superfície de fratura similares, o que pode estar associado ao valor limite de solubilidade. Em relação à composição 90/10 os pinos 70/30 e 50/50 apresentaram maior quantidade de poros em toda a amostra com uma morfologia esférica. Esse fato também foi observado por DSC onde não houve alteração significativa adicionando 30 e 50% do PCL-T.

A presença dos poros foi atribuída ao efeito plastificante do PCL-T, pois este foi o único diferencial na confecção dos pinos para origem dos poros.

Luciano[16] em seu estudo com membranas obtidas por evaporação de solvente observou que a adição do plastificante tri-etil-citrato na membrana densa tornava-a porosa. Segundo o autor o plastificante facilita a mobilidade das cadeias que se organizam ao redor de núcleos de cristalização já presente no polímero (semicristalino) e ocorre o surgimento de glóbulos. O aglomerado desses glóbulos gera uma estrutura porosa. Duarte[17] ao estudar de PLDLA/PCL-T na forma de membranas já havia verificado que a presença do PCL-T modificava a membrana de PLDLA tornando-a porosa.

Outra possível explicação para formação desses poros, baseado nos resultados, é que em função da presença de PCL-T provocar uma maior mobilidade nas cadeias, e estas por sua vez, na tentativa de se organizarem na nova condição que se encontram permitem a geração de pequenos espaços na estrutura, já que a distribuição aleatória das unidades levógiro e dextrógero, dificulta o processo de cristalização.

3.3 Ensaio de flexão a 3 pontos

O ensaio de flexão a 3 pontos mostrou que os pinos PLDLA/PCL-T obtidos por fusão têm características de um material dúctil.

Na Figura 4, pode se observar que a presença do PCL-T modificou o módulo de elasticidade do material. Essa mudança do módulo de elasticidade pode representar um ganho em possibilidade de aplicações da blenda.

A Tabela 2 mostra os dados obtidos do gráfico.

3.4 Análise in vitro dos pinos de PLDLA/PCL-T

O ensaio de flexão a 3 pontos mostrou como os pinos PLDLA/PCL-T se comportam num ambiente que simula as condições corpóreas, e a Tabela 3, apresenta esses dados.

Analisando a Tabela 3, pode se verificar que diferentemente do PLDLA puro, que durante o processo de degradação apresenta um aumento no modulo de elasticidade, para as demais situações em que temos as blendas, ocorre uma queda desse modulo de elasticidade no transcorrer dos dias, ainda que seja verificado em algumas composições, um aumento inicial deste parâmetro. Essa situação, de aumento do módulo de elasticidade no inicio do processo de degradação, é observada em inúmeros trabalhos[18-20], o que pode estar associado, por exemplo, a mudanças na massa molar do material.

Figura 2. Termogramas de DSC para o sistema PLDLA/PCL-T obtidas por fusão. (a) primeiro aquecimento; (b) segundo aquecimento.

Polímeros, 26(2), 185-191, 2016 189

Caracterização de pinos da blenda poli(L-co-D,L ácido láctico)/poli(caprolactona triol) (PLDLA/PCL-T) e análise das propriedade mecânicas dos pinos durante degradação in vitro

Figura 3. Micrografias de MEV da superfície (a) e de fratura (b) dos pinos PLDLA/PCL-T (100/0); da superficie (c) e de (d) fratura PLDLA/PCL-triol (90/10); da superfície (e) e de fratura (f) PLDLA/PCL-triol (70/30); da superfície (g) e de fratura (h) PLDLA/PCL-triol (50/50).

Duarte, M. A. T., Motta, A. C., & Duek, E. A. R.

Polímeros, 26(2), 185-191, 2016190

Foi verificado que quanto maior a quantidade de PCL-T na blenda, menor era o módulo elástico. Para o tempo de 24 semanas, as blendas, nas diferentes composições, apresentavam- se incapazes de serem submetidas ao ensaio mecânico, evidenciando o processo de degradação que se encontravam. O perfil de degradação que cada composição da blenda apresenta, acompanhado pelos dados. É importante salientar que a escolha de uma blenda, para uma determinada aplicação dependerá de sua capacidade em manter suas propriedades por um tempo determinado.

Tabela 3. Dados obtidos do ensaio de flexão nas diversas composições do PLDLA/PCL-triol nos tempos de estudo.

PLDLA/PCL-T(100/0)

Força

Máxima

(N)

Módulo de

Elasticidade (E)

(MPa)

Tensão máxima

(MPa)

Alongamento

na ruptura

(%)

0 semana 11,8 (±0,1) 3133 (±200) 50 (±2) 0,012 (±0,003)4 semanas 12,6 (± 0,4) 3462 (± 843) 67,8 (± 1,8) 0,02 (± 0,004)8 semanas 11,5 (± 0,6) 3791 (± 214) 59,5 (± 5,0) 0,016 (± 0,002)12 semanas 11,4 (± 0,3) 4024 (±428) 72,3 (± 5,7) 0,018 (± 0,002)16 semanas 9,9 (± 0,9) 3049 (± 285) 48,7 (± 9,8) 0,015 (± 0,002)24 semanas 7,8 (± 1,1) 3582 (± 367) 38,6 (± 4,9) 0,010 (± 0,001)PLDLA/

PCL-T (90/10)0 semana 8.9 (±0.2) 3212 (±200) 47,9 (±0,01) 0,016 (±0,0007)4 semanas 7,9 (±0,4) 3310(±238) 49,3 (±5,0) 0,06 (±0,079)8 semanas 6,9 (±0.1) 3293 (±111) 44 (±3) 0,013 (±0,001)12 semanas 4,2 (±0.8) 2046 (±552) 26,59 (±6,47) 0,013 (±0,003)16 semanas 3,7 (±0,2) 2773 (±431) 20,45 (±0,97) 0,007 (±0,001)24 semanas * * * *PLDLA/

PCL-T (70/30)0 semana 7,0 (±0,5) 2525 (±290) 43,4 (±3,0) 0,017 (±0,001)4 semanas 6,0 (±0,2) 2239 (±204) 38,3 (±1,5) 0,017 (±0,0004)8 semanas 5,9 (±0,5) 2894 (±398) 37,6 (±5,3) 0,013 (±0,00004)12 semanas 4,6 (±0,5) 2452 (±911) 20,9 (±5,0) 0,0089 (±0,002)16 semanas 3,2 (±1,1) 2149 (±586) 14,4 (±4,7) 0,006 (±0,0004)24 semanas * * * *PLDLA/

PCL-T(50/50)0 semana 4,4 (±0,9) 1738 (±132) 26 (± 4) 0,0149 (± 0,001)4 semanas 5,5 (±0,3) 2223 (±118) 37 (± 4) 0,017 (± 0,001)8 semanas 5,3 (±0,3) 2246 (±247) 31,1 (±2,8) 0,014 (± 0,002)12 semanas 4,3(±0,7) 2256 (±266) 22 (±8) 0,0092 (± 0,002)16 semanas 3,2 (±0,4) 1688 (±269) 14,2 (±1,4) 0,008 (± 0,001)24 semanas * * * *

*situação onde não foi possível fazer análise.

3.5 Análise macroscópica dos pinos nos diferentes tempos

A avaliação macroscópica da degradação dos pinos em função do tempo de imersão em tampão fosfato, para as diferentes composições de PLDLA/PCL-T, esta apresentada na Figura 5.

O PLDLA é um material amorfo e apresenta como característica ser transparente, conforme a composição 100/0 PLDLA, Figura 5a. Com a adição do PCL-T o material se torna levemente opaco, sendo tal opacidade bem visível na composição 50/50, Figura 5d. Após 24 semanas de degradação, verificou-se uma aparência completamente esbranquiçada

Figura 4. Curvas de força x deformação das blendas PLDLA/PCL-T.

Tabela 2. Valores de força máxima, módulo de elasticidade, tensão máxima e alongamento obtidos a partir do ensaio de flexão para os pinos PLDLA/PCL-T.

Composição: PLDLA/PCL-T

Força

Máxima

(N)

Módulo de

Elasticidade (E)

(MPa)

Tensão

(MPa)Alongamento

100/0 11,8 (±0,1) 3133 (±200) 49,9 (±2,6) 0,012 (±0,003)90/10 8.9 (±0,2) 3212 (±200) 47,9 (±0,01) 0,016 (±0,0007)70/30 7,0 (±0,5) 2525 (±290) 43,4 (±3,0) 0,017 (±0,001)50/50 4,4 (±0,9) 1738 (±132) 26 (± 4) 0,0149 (± 0,001)

Figura 5. Avaliação macroscópica dos pinos antes e após degradação (0 semana e 24 semanas).

Polímeros, 26(2), 185-191, 2016 191

Caracterização de pinos da blenda poli(L-co-D,L ácido láctico)/poli(caprolactona triol) (PLDLA/PCL-T) e análise das propriedade mecânicas dos pinos durante degradação in vitro

para o PLDLA, Figura 5e, o que pode ser atribuído à recristalização do material. Em todas as composições era nítida a constatação do processo de degradação, transcorridos 24 semanas, no entanto, a composição que se apresentava com um processo de degradação mais intenso, apresentando-se até na forma de fragmentos do pino, foi a PLDLA/PCL-T 90/10, Figura 5f. justificando, portanto a impossibilidade notada de se fazer o ensaio mecânico, conforme Tabela 3.

4. Conclusão

Baseado nos resultados apresentados pôde-se verificar a influência da concentração do PCL-T nos pinos de PLDLA, o qual modificou suas propriedades mecânicas, térmicas, morfológicas. A análise de DSC mostrou que as Tg do PLDLA diminuem na presença do PCL-T blendas, e que essa diminuição se mostrou significativa com 10% de PCL-T, não sendo notada uma queda mais expressiva desse patamar para concentrações maiores de PCL-T na blenda. A análise de MEV mostrou o aparecimento de poros no PLDLA quando o PCL-T estava presente enquanto que o teste de flexão mostrou uma diminuição do módulo de elasticidade, quando a concentração de PCl-T aumentava na blenda Esse aumento na flexibilidade dos pinos PLDLA, causado pelo PCL-T pode representar um aumento do numero de aplicações do material, em especial na área de engenharia tecidual.

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Enviado: Set. 12, 2014 Revisado: Ago. 28, 2015

Aceito: Set. 18, 2015

Polím

eros VOLUM

E XXVI - N° 2 - ABR/M

AIO/JUN - 2016

An_1Pa�g_SGS_Polimeros_Tintas_18x26,5cm.indd 1 5/31/16 4:36 PM